The interrelations between supermolecular structure,

properties and degradability of isotactic poly(1-butene)

Lubomır Benıcek

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Lubomır Benıcek. The interrelations between supermolecular structure, properties and degrad-ability of isotactic poly(1-butene). Materiaux. Universite Blaise Pascal - Clermont-Ferrand II,2009. Francais. <NNT : 2009CLF21921>. <tel-00725695>

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N° d’Ordre : D.U. 1921

UNIVERSITE BLAISE PASCAL U.F.R. Sciences et Technologies

ECOLE DOCTORALE DES SCIENCES FONDAMENTALES N° 605

THESE présentée pour obtenir le grade de

DOCTEUR D’UNIVERSITE

Spécialité : Chimie - Physique

Par BENÍČEK Lubomír Master en chimie et technologie des matériaux

THE INTERRELATIONS BETWEEN SUPERMOLECULAR STRUCTURE, PROPERTIES AND DEGRADABILITY OF ISOTACTIC POLY(1-BUTENE)

Soutenue publiquement le 27 avril 2009 devant la commission d’examen.

Président : Mme. Sophie COMMEREUC

Examinateurs : Mme. Berenika HAUSNEROVÁ (rapporteur) M. Philippe CASSAGNAU (rapporteur) M. Roman ČERMÁK M. Michel DUMON M. Martin OBADAL M. Miroslav RAAB (rapporteur) M. Vincent VERNEY

TOMAS BATA UNIVERSITY IN ZLÍN Faculty of Technology

Department of Polymer Engineering

DOCTORAL THESIS

Lubomír Beníček

THE INTERRELATIONS BETWEEN SUPERMOLECULAR STRUCTURE, PROPERTIES AND DEGRADABILITY OF ISOTACTIC POLY(1-BUTENE)

Programme: Chemistry and Materials Technology

Course: Technology of Macromolecular Substances

Co-supervisors: Dr. Vincent Verney doc. Ing. Martin Obadal, Ph.D.

Consultant: doc. Ing. Roman Čermák, Ph.D.

Year: 2009

3

ACKNOWLEDGEMENT In the first place, I would like to sincerely thank my supervisors Martin OBADAL and Vincent

VERNEY for giving me the opportunity of co-supervised Ph.D. study, valuable advices and inspirational

suggestions that have helped me to successfully complete this thesis. I also thank Roman ČERMÁK for his

help, ideas and discussion during my research as my consultant.

Further, I would like to thank Sophie COMMEREUC, Berenika HAUSNEROVÁ,

Philippe CASSAGNAU, Michel DUMON and Miroslav RAAB for being jury members and for their

valuable comments.

Special thank belongs:

French government for the financial support (BGF) and Marie-Claire COMTE-ROME from CROUS

Clermont-Ferrand.

Faculty of Technology, Tomas Bata University in Zlín and Laboratoire de Photochimie Moleculaire

et Macromoleculaire, Université Blaise Pascal for technical and scientific support with experiments.

My colleagues, friends and students who helped me with experiments in the both institutions where

I have been studying: Lenka, Solène, Jana, Kristýna, Slavomíra, Martina, Monika, Branislav, Harout,

Michal K., Denis, Matthieu, Michal M., Tomáš, Guillaume, Joël CELLIER, Lawrence FREZET and

Milan ŽALUDEK.

Last but not least to my fiancée Libuše, parents and family for their great support, patience and love.

4

RÉSUMÉ Cette thèse porte sur l’inter relation entre la superstructure moléculaire, les propriétés et la

dégradabilité du poly butène isotactique. L’évolution la plus importante est une transformation de phase

particulière de la forme cristalline II à la forme cristalline I après cristallisation du polymère à partir de l’état

fondu. Cette transformation a été caractérisée par calorimétrie, par diffraction des rayons X aux grands

angles et par l’analyse des propriétés mécaniques en fonction des conditions de thermique de recuit. L’effet

de la mise en forme sur la superstructure résultante a été mis en évidence sur des échantillons obtenus par

injection ou par extrusion en utilisation les propriétés mécaniques en traction. Du fait de la très grande

sensibilité du PB-1 vis-à-vis de la photodégradation des effets significatifs sur les propriétés thermiques et

sur la transformation de phase ont été observés. Une caractérisation multi-échelle en photodégradation a été

réalisée à partir d’analyses calorimétriques, d’analyse en spectroscopie infrarouge, de la diffraction des

rayons X et des propriétés mécaniques en traction et ceci tant en photovieillissement en conditions accélérées

qu’en conditions d’exposition naturelle.

Mots clés: poly(1-butène) isotactique, transformation de phase, structure, propriétés,

photodégradation

ABSTRAKT

Tato dizertační práce je zaměřena na studium vzájemných vztahů mezi nadmolekulární strukturou,

vlastnostmi a degradabilitou izotaktického poly(1-butenu) (PB-1). Nejvýznamnější vývoj v PB-1 je jeho

zvláštní fázová transformace z formy II do formy I, jenž nastává po krystalizaci z taveniny. Tato fázová

transformace je charakterizována pomocí kalorimetrického měření, široko-úhlé rentgenografie a

mechanickými zkouškami, kde byly pozorovány změny v závislosti na podmínkách stárnutí. Vliv

zpracovatelské technologie – vstřikování a vytlačování - na výslednou nadmolekulární strukturu byl

prokázán na základě mechanických zkoušek. Jelikož je PB-1 velmi náchylný k fotodegradaci, významné

změny byly zaznamenány v termálních vlastnostech a fázové transformaci. Široce zaměřená charakterizace

důsledků fotodegradace byla provedena pomocí kalorimetrických měření, infračervené spektroskopie,

reologie, rentgenografie, mikroskopických technik a tahových zkoušek. Rovněž byl posouzen vliv

rozdílných degradačních procesů – venkovní stárnutí a zrychlené UV ozařování – které vedly k odlišným

výsledkům ovlivňující výsledné makroskopické vlastnosti.

Klíčová slova: izotaktický poly(1-buten), fázová transformace, struktura, vlastnosti, fotodegradace

5

ABSTRACT This dissertation work is focused on the interrelations between supermolecular structure, properties

and degradability of isotactic poly(1-butene) (PB-1). The most important evolution in PB-1 is a peculiar

phase transformation from form II to form I after crystallization from melt. This phase transformation is

characterized using calorimetry, wide-angle X-ray scattering and mechanical testing where changes in

dependence on annealing conditions in time are observed. The effect of processing on resulting

supermolecular structure is proven on injection-molded and extruded specimens using tensile testing.

Because of high sensitivity of PB-1 to photodegradation, significant effects on thermal properties and the

phase transformation is observed. Multiscale characterization is performed using calorimetric measurements,

infra-red spectroscopy, rheology, X-ray scattering, microscopic techniques and tensile testing to describe

consequences of photodegradation. As well, the effect of various degradation processes – natural weathering

and accelerated UV irradiation lead to different results affecting resulting macroscopic properties.

Keywords: isotactic poly(1-butene), phase transformation, structure, properties, photodegradation

6

TABLE OF CONTENTS

ACKNOWLEDGEMENT .................................................................................................................... 3 

RÉSUMÉ .............................................................................................................................................. 4 

ABSTRACT ......................................................................................................................................... 5 

TABLE OF CONTENTS ..................................................................................................................... 6 

LIST OF ABBREVATIONS ................................................................................................................ 9 

LIST OF FIGURES ............................................................................................................................ 11 

1  INTRODUCTION ....................................................................................................................... 14 

2  BIBLIOGRAPHIC STUDY ........................................................................................................ 15 

2.1  Isotactic poly(1-butene) ....................................................................................................... 15 

2.1.1  History ............................................................................................................................ 15 

2.1.2  Molecular structure ......................................................................................................... 16 

2.1.3  Supermolecular structure ................................................................................................ 16 

2.1.4  Crystallization ................................................................................................................. 17 

2.1.5  Polymorphic transformation ........................................................................................... 18 

2.2  Properties of PB-1 ............................................................................................................... 23 

2.2.1  Mechanical properties ..................................................................................................... 23 

2.2.2  Physical properties .......................................................................................................... 24 

2.2.3  Processing and application ............................................................................................. 25 

2.3  Photodegradation ................................................................................................................. 27 

2.3.1  Radical oxidation ............................................................................................................ 28 

2.3.2  Aspects of natural weathering......................................................................................... 28 

2.3.3  Crystallization of degraded polyolefins .......................................................................... 29 

3  SUMMARY AND AIMS ............................................................................................................ 31 

4  MATERIALS AND EXPERIMENTAL TECHNIQUES ........................................................... 32 

4.1  Materials .............................................................................................................................. 32 

4.2  Sample preparation .............................................................................................................. 32 

4.2.1  Compression molding ..................................................................................................... 32 

7

4.2.2  Injection molding ............................................................................................................ 32 

4.2.3  Extrusion ......................................................................................................................... 33 

4.3  Ageing ................................................................................................................................. 33 

4.3.1  Annealing ........................................................................................................................ 33 

4.3.2  Natural weathering .......................................................................................................... 33 

4.3.3  UV irradiation ................................................................................................................. 33 

4.4  Testing and characterization ................................................................................................ 35 

4.4.1  Tensile testing ................................................................................................................. 35 

4.4.2  Polarized light microscopy ............................................................................................. 35 

4.4.3  Stereomicroscopy ........................................................................................................... 35 

4.4.4  Scanning electron microscopy ........................................................................................ 35 

4.4.5  Wide-angle X-ray scattering ........................................................................................... 35 

4.4.6  Melt rheology.................................................................................................................. 37 

4.4.7  Infrared spectroscopy ...................................................................................................... 39 

4.4.8  Differential scanning calorimetry ................................................................................... 39 

5  STUDY OF TRANSFORMATION IN PB-1.............................................................................. 41 

5.1  DSC study ............................................................................................................................ 41 

5.2  Evolution of properties of injection-molded PB-1 .............................................................. 45 

5.2.1  Tensile modulus .............................................................................................................. 45 

5.2.2  Tensile yield strength ...................................................................................................... 47 

5.2.3  Tensile elongation at break ............................................................................................. 48 

5.2.4  Tensile strength at break ................................................................................................. 50 

5.3  Evolution of structure and properties in PB-1 extrudates .................................................... 51 

5.3.1  WAXS evolution ............................................................................................................ 51 

5.3.2  Polarized light microscopy ............................................................................................. 53 

5.3.3  Tensile modulus .............................................................................................................. 54 

5.3.4  Tensile elongation at break ............................................................................................. 56 

5.3.5  Tensile strength at break ................................................................................................. 58 

6  MULTISCALE CHARACTERIZATION OF DEGRADATION IN PB-1 ................................ 61 

8

6.1  Early stages .......................................................................................................................... 61 

6.1.1  DSC study ....................................................................................................................... 61 

6.1.2  FTIR characterization ..................................................................................................... 67 

6.1.3  Rheology ......................................................................................................................... 70 

6.2  Prolonged stages .................................................................................................................. 73 

6.2.1  DSC study ....................................................................................................................... 73 

6.2.2  FTIR-ATR characterization ............................................................................................ 73 

6.2.3  Rheology ......................................................................................................................... 74 

6.2.4  WAXS characterization .................................................................................................. 77 

6.2.5  Stereomicroscopy ........................................................................................................... 78 

6.2.6  Scanning Electron Microscopy ....................................................................................... 78 

6.3  Degradation of extruded PB-1 ............................................................................................. 81 

6.3.1  Mechanical properties ..................................................................................................... 81 

6.3.2  FTIR-ATR characterization ............................................................................................ 83 

6.3.3  WAXS characterization .................................................................................................. 85 

6.3.4  Stereomicroscopy ........................................................................................................... 87 

6.3.5  Effect of temperature on crystallinity and form II content ............................................. 88 

7  CONCLUSION AND PERSPECTIVES ..................................................................................... 90 

8  REFERENCES ............................................................................................................................ 92 

AUTHOR’S PUBLICATIONS .......................................................................................................... 96 

CURRICULUM VITAE ..................................................................................................................... 98 

9

LIST OF ABBREVATIONS Å Angstrom

C-M Compression-molded

cx Relative crystallinity

DSC Differential scanning calorimetry

FTIR Fourier Transform Infrared

FTIR-ATR Fourier Transform Infrared - Attenuated Total Reflectance

G* Complex dynamic modulus

G’ Storage modulus

G” Loss modulus

h Parameter of the relaxation-time distribution

HDPE High-density polyethylene

Hg Mercury

I Total intensity

Ic Crystalline part

IR Infrared

m1, m2 First and second melting corresponding to form I and II

mtr Melting after transformation is completed

MW Molecular weight

MWD Molecular weight distribution

NMR Nuclear magnetic resonance

NOx Nitrogen oxide

PB-1 Isotactic poly(1-butene)

PE Polyethylene

PLM Polarized light microscopy

PP Isotactic polypropylene

Rg Radius of gyration

SAXS Small-angle X-ray scattering

SEM Scanning electron microscopy

tan δ Phase angle

Tc Crystallization temperature

TEM Transmission electron microscopy

Tm Melting temperature

UHMWPE Ultra-high molecular weight polyethylene

UV Ultra-violet

WAXS Wide-angle X-ray scattering

Δl Elongation

η* Complex viscosity

η’, η” Real and imaginary viscosity components

10

η0 Zero shear viscosity

λ0 Average relaxation time

ω Frequency

11

LIST OF FIGURES Fig. 1 Chemical formula of monomer and PB-1 polymer ............................................................................... 16 

Fig. 2 Interconversion between various physical forms of PB-1. (For simplicity, only an amyl acetate

solution is considered in this figure) [32] ........................................................................................... 19 

Fig. 3 Dependence of phase transformation of form II to form I on ambient temperature and pressure [34] . 19 

Fig. 4 Tensile behavior of PB-1 vs. other polyolefins [34] ............................................................................. 24 

Fig. 5 Scheme of standard mechanism of the oxidation in polymer chain [62] .............................................. 28 

Fig. 6 Scheme of a SEPAP 12-24 .................................................................................................................... 34 

Fig. 7 Spectra of lamps vapor Hg medium pressure used in a Sepap 12-24 ................................................... 34 

Fig. 8 Wide-angle X-ray scattering patterns at different times [26] ................................................................ 37 

Fig. 9 Deconvolution of the DSC curve obtained by PEAK-FIT V4 [105] .................................................... 40 

Fig. 10 Evolution of form I content in DP 0401M for various annealing temperatures .................................. 42 

Fig. 11 Evolution of form I content in PB 0300M for various annealing temperatures .................................. 42 

Fig. 12 Evolution of form I content in PB 8640M for various annealing temperatures .................................. 43 

Fig. 13 Evolution of form I content in the whole temperature range for all materials in selected annealing

days ..................................................................................................................................................... 44 

Fig. 14 Evolution of tensile modulus in injection-molded PB 0300M ............................................................ 46 

Fig. 15 Evolution of tensile modulus in injection-molded DP 0401M ............................................................ 46 

Fig. 16 Evolution of tensile modulus in injection-molded PB 8640M ............................................................ 47 

Fig. 17 Evolution of tensile strength at yield in injection-molded DP 0401M ................................................ 47 

Fig. 18 Evolution of tensile elongation at break in injection-molded PB 0300M ........................................... 48 

Fig. 19 Evolution of tensile elongation at break in injection-molded PB 8640M ........................................... 49 

Fig. 20 Evolution of tensile strength at break in injection-molded PB 0300M ............................................... 50 

Fig. 21 Evolution of tensile strength at break in injection-molded PB 8640M ............................................... 50 

Fig. 22 WAXS evolution of the form I content in extruded PB 0300M for various annealing temperatures

upon time ............................................................................................................................................ 52 

Fig. 23 Images of morphology in extruded PB-1(bottom is border of specimen) ........................................... 53 

Fig. 24 Evolution of tensile modulus in extruded PB 0300M ......................................................................... 55 

Fig. 25 Evolution of tensile modulus in extruded DP 0401M ......................................................................... 55 

Fig. 26 Evolution of tensile modulus in extruded PB 8640M ......................................................................... 56 

Fig. 27 Evolution of tensile elongation at break in extruded PB 0300M ........................................................ 57 

Fig. 28 Evolution of tensile elongation at break in extruded DP 0401M ........................................................ 57 

Fig. 29 Evolution of tensile elongation at break in extruded PB 8640M ........................................................ 58 

Fig. 30 Evolution of tensile strength at break in extruded PB 0300M ............................................................ 59 

Fig. 31 Evolution of tensile strength at break in extruded DP 0401M ............................................................ 59 

Fig. 32 Evolution of tensile strength at break in extruded PB 8640M ............................................................ 60 

12

Fig. 33 Melting behavior of PB 0300M and DP 0401M (m1 – first melting, m2 – second melting, mtr –

melting after transformation is completed) ......................................................................................... 62 

Fig. 34 Crystallization behavior of PB 0300M and DP 0401M ...................................................................... 62 

Fig. 35 Melting and crystallization behavior of DP 0401M exposed to natural weathering, (m1 – first

melting, m2 – second melting, mtr – melting after transformation is completed) ............................. 63 

Fig. 36 Evolution of form I content during transformation of degraded specimens for various irradiation

times in DP 0401M ............................................................................................................................ 64 

Fig. 37 Evolution of ratio of form I to II during transformation of degraded specimens for various irradiation

times in DP 0401M ............................................................................................................................. 65 

Fig. 38 Evolution of form I content during transformation of degraded specimens for various irradiation

times in PB 0300M ............................................................................................................................ 65 

Fig. 39 Evolution of ratio of form I to II during transformation of degraded specimens for various irradiation

times in PB 0300M ............................................................................................................................. 66 

Fig. 40 Evolution of form I content during transformation of degraded specimens for various exposition

times in DP 0401M ............................................................................................................................ 66 

Fig. 41 Evolution of ratio of form I to II during transformation of degraded specimens for various exposition

times in DP 0401M specimens ........................................................................................................... 67 

Fig. 42 FTIR spectra of DP 0401M - evolution in carbonyl area during UV irradition .................................. 68 

Fig. 43 FTIR spectra of PB 0300M - evolution in carbonyl area during UV irradiation ................................ 68 

Fig. 44 FTIR spectra of DP 0401M - evolution in carbonyl area during natural exposition ........................... 69 

Fig. 45 Comparison of absorbancies at 1712 cm-1 in UV irradiated and naturally exposed specimens .......... 70 

Fig. 46 Cole-Cole plot of DP 0401M up to 10 hours of UV irradiation .......................................................... 71 

Fig. 47 Cole-Cole plot of PB 0300M up to 10 hours of UV irradiation .......................................................... 71 

Fig. 48 Evolution of zero shear viscosity and distribution parameter during UV irradiation ......................... 72 

Fig. 49 Melting and crystallization behavior of PB 0401M (m1 – first melting, m2 – second melting, mtr –

melting after transformation is completed) ......................................................................................... 73 

Fig. 50 FTIR-ATR spectra of DP 0401M - evolution of carbonyl area during UV irradiation ....................... 74 

Fig. 51 Cole-Cole plot of DP 0401M up to 20 hours of UV irradiation .......................................................... 75 

Fig. 52 Cole-Cole plot of PB 0300M up to 20 hours of UV irradiation .......................................................... 75 

Fig. 53 Evolution of zero shear viscosity and distribution parameter during UV irradiation in DP 0401M ... 76 

Fig. 54 Evolution of zero shear viscosity and distribution parameter during UV irradiation in PB 0300M ... 76 

Fig. 55 Evolution of crystallinity and form II content in UV irradiated specimens DP 0401M ..................... 77 

Fig. 56 Evolution of crystallinity and form II content in naturally exposed specimens DP 0401M ............... 77 

Fig. 57 Images of surface degradation in DP 0401M in various UV irradiation times ................................... 78 

Fig. 58 SEM images of surface degradation in the materials for various UV irradiation times ...................... 79 

Fig. 59 SEM images of surface degradation in the DP 0401 naturally degraded at various exposition times 80 

Fig. 60 Evolution of tensile modulus in UV irradiated and naturally exposed specimens .............................. 81 

13

Fig. 61 Evolution of tensile stress at break in UV irradiated and naturally exposed specimens ..................... 82 

Fig. 62 Evolution of tensile elongation at break in UV irradiated and naturally exposed specimens ............. 82 

Fig. 63 FTIR-ATR spectra of UV irradiated specimens.................................................................................. 84 

Fig. 64 FTIR-ATR spectra of naturally exposed specimens ........................................................................... 84 

Fig. 65 Comparison of absorbancies at 1712 cm-1 in UV irradiated and naturally exposed specimens .......... 85 

Fig. 66 Evolution of crystallinity and form II content in UV irradiated specimens ........................................ 86 

Fig. 67 Evolution of crystallinity and form II content in naturally exposed specimens .................................. 86 

Fig. 68 Images of extruded tapes DP 0401M in various UV irradiation times ............................................... 87 

Fig. 69 Evolution of crystallinity in annealed and degraded specimens ......................................................... 88 

Fig. 70 Evolution of form II content in annealed and degraded specimens .................................................... 89 

14

1 INTRODUCTION

This Ph.D. thesis (these en co-tutelle - co-supervised dissertation) has arisen from previous cooperation between Department of Polymer Engineering, Faculty of technology, Tomas Bata University in Zlín and Laboratoire de Photochimie Moleculaire et Macromoleculaire, Université Blaise Pascal, Clermont-Ferrand. The French government (Bourse de Gouvernement Francais) financed 18 months long study in France after successful presentation of the proposed project. The realization was performed during last 3 years with alternation of 6 months in France and 6 months in Czech Republic.

Studies related to the phase transformation of isotactic poly(1-butene) (PB-1) possesses, in general, rather high practical importance, as this phenomenon directly influences all material (product) properties; surprisingly, there are virtually no articles screening overall consequences arising from and after processing and during lifetime. Actually, these difficult interrelations are, in fact, a principal reason, why PB-1 has not found wide spectrum of applications yet, although its properties are favorably placed between isotactic polypropylene (PP) and high-density polyethylene (HDPE) and in some of them, even it exceeds their properties, e.g. creep behavior. Many articles have been devoted to explain a mechanism of the phase transformation, however these do not fully answer industrial questions how much time it takes until phase transformation is completed and at which temperature the processed products should be stored to obtain the fastest transformation rate. Both questions are closely connected with economical aspects of PB-1 production. A PB-1 degradability has not been practically studied at all, in term of assessment of durability the commercially produced PB-1. In addition, the question of subsequent recycling is nowadays frequently discussed; the properties of recycled PB-1 are practically unknown, neither the effect of oxidation on the phase transformation in such degraded material.

Thus, the knowledge of the interrelations between the supermolecular structure and the properties is therefore an important prerequisite enabling an estimation of the applicability of polymer materials. For this reason and mentioned questions, thesis is basically divided into two main parts.

The first part is concerned to the phase transformation of PB-1 itself; initially, using differential scanning calorimetry the transformation behavior is investigated in respect to different annealing condition. Afterward, the phase evolution is studied in both injection-molded and extruded PB-1 specimens using tensile testing with supplementary structural analysis being done with the help of X-ray scattering and microscopy.

The second part deals with a study of degradability of compression molded and extruded specimens under natural and accelerated UV conditions. Calorimetric characterization is performed to determine thermal properties and subsequently to observe the effect of degradation on phase transformation behavior. The particular attention is paid to a rheological characterization, which highly is sensitive to detect early stages of degradation. The property evolution on macroscopic scale of degraded specimens predicting its lifetime is observed through tensile behavior. The changes on supermolecular structural level are again screened using X-ray scattering, microscopy and infrared spectroscopy.

15

2 BIBLIOGRAPHIC STUDY

2.1 Isotactic poly(1-butene)

2.1.1 History

Isotactic poly(1-butene) (PB-1) was first synthesized in 1954, one year after polypropylene. It took

another 10 years until Chemische Werke HÜLS, Germany, started the first industrial production in 1964.

Vestolen BT was introduced to the market [1,2].

Independently, Mobil Oil in the USA developed their own PB-1 process technology and built a small

industrial plant in Taft, Louisiana, in 1968. In early 70’s the plant was taken over and operated by Witco

Chemical Corporation.

In 1973 HÜLS withdrew Vestolen BT from the market after some manufacturing issues in their

polymerization plant.

At the end of 1977 Shell Chemicals USA, a subsidiary of Shell Oil Company, acquired the PB-1

business from Witco, including the Taft plant. Shell then started a major investment program to improve the

product quality and to increase the production capacity.

In 1988, the management of the PB-1 portfolio was put into the hands of Shell’s subsidiary Montell

Polyolefins. Exactly two years later, in October 2000, Montell Polyolefins merged with Targor GmbH and

Elenac GmbH to form Basell Polyolefins.

In December 2000 Basell’s management board approved an investment to construct a new plant in

Europe. It utilizes a new process technology, which has been developed by Basell’s Research and

Development Centre in Ferrara, Italy. The new plant was constructed in Moerdijk, The Netherlands in 2003

[3]. In December 2007, Basell and Lyondell Chemical Company completed their merger to create

LyondellBasell – one of the world’s largest polymers, chemicals and fuels companies.

Apart from LyondellBasell, Mitsui Chemicals also manufacture PB-1 on their multi-purpose plants

in Japan.

16

2.1.2 Molecular structure

There are three possibilities of stereochemical arrangement of the polymeric chains in poly(1-

butene): syndiotactic, isotactic and atactic. While PB-1 was successfully synthesized and extensively studied

in the mid-1950s, the synthesis of syndiotactic poly(1-butene) was just achieved by the recent discovery of

the homogeneous metallocene catalytic systems [4].

Fig. 1 Chemical formula of monomer and PB-1 polymer

The stereochemical configuration or “relative handedness” of successive monomer units in

polyolefins such as PB-1 can strongly affect the physical and mechanical properties of the polymers. The

stereo-specific Ziegler-Natta catalyst system is employed and the polymerization conditions are used to

control polymer stereoregularity. The tacticity of PB-1 is strongly dependent on the catalyst system used,

polymerization temperature, and reaction diluents [5]. Several methods can be used to quantify the

isotacticity of PB-1. These methods include selective solvent extraction, infrared spectroscopy, scattering,

NMR spectroscopy [4,6].

2.1.3 Supermolecular structure

For over 50 years it has been known that crystalline PB-1 occurs in multiple supermolecular forms.

As all stereoregular polyolefins, PB-1 exhibits also polymorphic crystallization behavior. Five different

crystalline forms have been recently reported in the literature, which are referred to as I, II, III, I’, and II’ [7].

In 1955, Natta et al. [8,9] reported the first two forms, with one of these being sufficiently unstable

that it converted rapidly to the other, at pressures of around 10 MPa. The most common of these, called form

I, was extensively studied in 1960. The crystal structure of form II was determined 3 years later, in 1963.

Cojazzi [10] reported a form III of PB-1, obtained by evaporation from solution, in 1976.

Form II possesses 11/3 helical conformation with a tetragonal unit cell when crystallizing from the

melt [11], with identity period of 6.8 Å [12]. Form II is unstable and transforms into a stable 3/1 helix

conformation known as form I with twinned hexagonal [8] (rhombohedral) unit cells at room temperature

and atmospheric pressure [8], with identity period of 6.5 Å [12]. The form I is being also assigned as trigonal

form by some authors [13].

H

C

H

C

H CH2

CH3

H

C

H

C

H CH2

CH3

n

catalyst

heat & pressuren x

1-butene poly(1-butene)

17

Orthorombic form III with 4/1 helical conformation has been observed in films of PB-1 precipitated

from certain solvents [38].

The appearance of untwinned hexagonal form I’ and form II’ is related to crystallization under

pressure [14] or from solution in case of form I’ [15,16].

The first effort to study the kinetics of the crystallization and subsequent phase transformationhas

been offered by Boor and Mitchell [17].

2.1.4 Crystallization

Cooke et al. studied using WAXS, SAXS and DSC the changes in the Avrami exponent at different

crystallization temperatures, which are attributed to a change in nucleation mechanism [18].

The effect of molecular and processing parameters on shear-induced crystallization has been

preliminary studied by Baert et al. [19]. They expect that the growth rate at high Tc is determined by the

formation of stable nuclei and not so much by diffusion of polymer chains as in the case at low Tc, which

results from Lauritzen-Hoffman polymer crystallization theory [20]. They also reported occurrence of the

two different crystalline structures – besides expected spherulites, a square-like morphologies called hedrites

being earlier reported [21-23]. Fu et al. [23] explained this effect as a change in the crystallization

mechanism from folded-chain to chain-extended crystallization taking when the crystal thickness gets larger

than the radius of gyration Rg of the chains.

The morphology of melt-crystallized PB-1 depends remarkably on the temperature of crystallization.

When crystallizing the iPB-1 at temperatures lower than 60 °C, spherulitic structures were observed. Flat-on

lathlike crystals could be formed at temperatures higher than 90 °C. In the temperature range from 60 to

90 °C, the spherulitic and flat-on lathlike structures coexist in the same sample [24].

The application of SCORIM (shear controlled orientation in injection molding) in PB-1 was studied

by Kalay et al. [4]; a shish-kebab morphology was formed by a SCORIM injection molding, improved

mechanical properties were observed.

The correlation between the gel time (measured from the change in tg δ) and quiescent/quasi-

quiescent crystallization onset time of PB-1 using rheological measurement was studied by Hadinata et al.

[25].

18

2.1.5 Polymorphic transformation

The most important phenomenon is following peculiar transformation - upon solidification from the

melt PB-1 crystallizes into form II which is kinetically favored. In this stage iPB-1 is a rather soft,

mechanically weak material. Over several days the material hardens by transforming into form I, which is

thermodynamically favored. This transformation results in the desirable properties of the material [12,26,27].

The melting point increases from 120 to 135 °C.

The kinetics of this phase transformation is known to be influenced by strain [28], pressure [29], and

temperature [27] (see Fig. 3). Actually, this rather slow kinetics is the principal reason why the commercial

development of PB-1 has been much delayed.

Nakafuku and Miyaki [29] studied the effect of pressure on the crystallization behavior of PB-1 and

reported that its melt crystallization under high pressure produces stable form I’, which shows the same

X-ray diffraction pattern as form I but has a much lower melting temperature (96 versus l30 °C) at

atmospheric pressure. Above 200 MPa, form I’ and form II’ are crystallized from the melt. Form II’ shows

the same X-ray diffraction pattern as form II, but a lower melting temperature than form II. Form II’ is

metastable at atmospheric pressure and transforms to form I’ on standing at room temperature [4].

Maring et al. [6,30] extensively studied a dynamic in the crystalline polymorphic forms I, II and III

using NMR. Later, Miyoshi et al. [31] elucidated that side-chain conformation in form II is disordered as

well as in amorphous sample, and the side-chain mobility in form II is higher than that in the amorphous

sample. Furthermore, the distributed main-chain conformations in form II are partially ordered with

reduction of the side-chain mobility. They suggested that the side-chain dynamics influences the main-chain

conformation in form II.

The samples as-polymerized with the Ziegler-type catalysts contain form I’ and form III. Melt

crystallization usually leads to the formation of form II. However, when PB-1 is melt-crystallized under high

hydrostatic pressure (>100 MPa), by a special meltstretching technique or onto specific substrates for

epitaxial growth, formation of form I’ has been reported. The formation of form III has also been reported

when PB-1 is melt-crystallized with a specific nucleating agent. In the case of crystallization from a dilute

amyl acetate solution, all three forms can be obtained by controlling dissolution and crystallization

temperatures. When form I’ and III were formed directly from the melt or from the amyl acetate solution,

they transform into form II by annealing at 90 ~ 100 °C (probably through a melting-recrystallization

process) [32]. The crystallization and morphology of the form III has been studied by Lotz and Thierry [33].

The interconversion between various physical forms of PB-1 is schematically shown in Fig. 2 [32].

19

Fig. 2 Interconversion between various physical forms of PB-1. (For simplicity, only an amyl acetate solution is considered in this figure) [32]

Fig. 3 Dependence of phase transformation of form II to form I on ambient temperature and pressure [34]

monomer melt solution

form III

form I’

form II

form I

nucleatingagent

polymerization

annealing 90 – 100 °C

ageing at room temperature

annealing 90 – 100 °C

annealing88 – 98.5 °C

below 100 MPa

epitaxial crystallizationabove 100 MPa

melt‐stretch

-40 -20 0 20 40 60 80 100 120 140 160

-3

-2

-1

0

1

0.1 MPa 50 MPa

120 MPa

150 MPa

log

reci

proc

al h

alf t

ime

t 1/2-1

(min

-1)

temperature (°C)

20

It is interesting that form II gradually transforms into form I during annealing at room temperature.

In the case of solution grown single crystals, the morphological change induced by this transformation is not

observed except for formation of small cracks. The form I that was obtained by the transformation from form

II is the most stable and does not change into form II even when the sample is annealed again. In this

context, the form I which is formed directly and can transform into form II by annealing is denoted as form

I’. Because form I’ showed broader crystalline reflection peaks in the X-ray spectrum than form I, form I’

has been regarded to be an imperfect form I with many defects. It is still unclear, in terms of the molecular

movements, how form II transforms into form I without changing the original shape [32].

In the early work by Miller and Holland [35], the formation of form I crystals twinning along the

diagonals of the original squareshaped single crystal of form II was proposed; it was assumed that each

quadrant of the square is an “untwinned” form I and adjacent quadrants are rotated by 90° with respect to

each other. However, the twinning may not be essential for the transformation.

Fujiwara [36] investigated the transformation of oriented form II specimens and found that the

“untwinned” form I appears by applying shear stress; the orientation of form I changed according to the

relative direction of the stress. He proposed a nucleation mechanism for the transformation.

Kopp et al. [12] investigated the form II single crystals grown from an octanol solution by

transmission electron microscopy (TEM), and their observation was explained by the molecular mechanism

proposed by Fujiwara [36]. Tosaka et al. [32] recognized that, the transformation of PB-1 from form II to

form I is initiated by stress, the nucleation of form I may occur as was proposed by Fujiwara [36]. However,

another type of nucleation mechanism, which creates the “twinned” form I seems to exist. Once the form I

domain is nucleated, it may grow by reeling in the molecular chains from the surrounding form II domains,

creating new stems for form I. The traction force to reel in the molecules may sometimes rotate the form I

domain itself. After the nucleation, the form I domain may be able to grow even when there is misalignment

in the relative orientation with the surrounding form II domain. There is no specific crystallographic

direction along which the form I domain tends to grow, and accordingly, the shape of the form I domain is

irregular [32].

Marigo et al. [26] studied phase transformation II to I by WAXS and SAXS and suppose that a

twofold mechanism of transformation could take place in PB-1. The transition nucleation seems to be

localized on lamellar distortion points, and the transition itself involves the rearrangement of lamellae and of

lamellar stacks. Moreover, a further crystallization of the amorphous phase into form I, seems to take place,

with the appearance of new thin lamellae inside lamellar stacks.

Kaszonyiova et al. [37] found that organic and inorganic nucleating agents commonly used for

polypropylene did not influence the polymorphism, but only crystallization rate.

21

It is reported that phase transformation is not accompanied by changes in macroscopic morphology

and the overall crystallinity [38]: at most, an increase of the degree of crystallinity of about 5 % for samples

quenched from melt.

While earlier studies of electron microscopy suggested that the phase transformation might produce

fragmentation of lamellae into small blocks separated cracks [39], recent highresolution transmission

electron microscopy studies could not provide an evidence of such cracks in transformed solution grown

lamellae [32].

Several articles dealt with direct crystallization of the form I from the melt in solutions [12,13,40-

42], however these attempts lead to single crystal without any perspective of industrial utilization.

Azzurri et al. [43,44] studied phase transformation by measurement of microindentation hardness

and calorimetry. They observed that the form I has notably higher hardness than the from II. A good

agreement was found between the degree of transformation evaluated from microhardness data and the

calorimetry measurements. It is suggested, also by Gohil et al. [45] that the transformation rate does not

depend on molecular mass (studied on PB-1 homopolymer supplied by Basell). High crystallization

temperatures increase the transformation rate. The increase in the rate of the process with decreasing fraction

of amorphous material is tentatively explained on the basis of the role by taut molecules at the crystal –

amorphous interphase eventually favoring the helix extension needed for the phase transformation.

On the contrary, former study of Chau et al. [46] claimed that the transformation rate is lower in the

high molecular weight sample and taut tie molecules cannot be a predominant factor which determines the

transformation rate.

Earlier study of Luongo and Salovey [47] discovered from their infrared spectroscopic study that the

rate of the phase transformation depends on the thickness of the compression molded films, decreasing with

increasing thickness.

Choi and White [48] observed that in thin melt spun fibers, transformation rate II – I increases with

the level of polymer chain orientation developed during melt spinning. In case of thick rods, the

transformation rate is faster in the core section and slow at the outer surface. It suggested to be due to the

development of quench stresses.

On the other hand, the effect of comonomeric units randomly distributed along the chain has been

less studied by Foglia [49], Turner-Jones [50] - investigation of copolymers of 1-butene with various

α-olefins. They observed that copolymerization of PB-1with α-olefins with less than 5 carbon atoms

increases the transformation rate while long linear α-olefins or branched co-units speed up the rate of the

process. Turner-Jones [50] explained these results by considering that the insertion of co-units having a small

cross-sectional area favors the 31 helical conformation and therefore increases the transformation rate. On the

22

contrary, long linear co-units, with more than 5 carbons, or branched comonomers, stabilize the 113 helix due

to steric factors and hinder the II - I transformation. Recently, Azzurri et al. [51] reinvestigated the phase

transformation in a series of PB-1 copolymers containing randomly distributed ethylene units. It was

observed that, similar to the homopolymers, the rate of the transformation does not depend on the average

length of constituent chains. On the other hand, a small amount of this type of co-unit is sufficient to strongly

accelerate the transformation to the extent that, at room temperature, only a few hours is sufficient to

complete the process, while in case of the homopolymers it takes several days.

It was found, in several studies that the presence of nanofillers in PB-1 matrices such as layered

silicates, nanotubes, nanoparticles, modified clays increases the transformation rate [52-57]. Wanjale and Jog

[52] suggested from their study on PB-1 with carbon nanotubes that it is attributed to the decrease in the

amorphous content and to the enhanced nucleation of form I due to disordered crystallite morphology.

23

2.2 Properties of PB-1 Isotactic poly(1-butene) can be regarded as a polymer having a number of attractive properties that

distinguish it from the most common polyolefins like isotactic polypropylene or polyethylene. PB-1 is a

semi-crystalline polymer with high isotacticity and consequently high crystallinity. Generally, the crystalline

part determines a number of intrinsic characteristics of the polymer. It exhibits advantages over the other

polyolefins in toughness, tear strength, flexibility, creep, stress cracking resistance, impact resistance and

abrasion resistance. The crystalline layers in PB-1 are connected by entangled tie-molecules. These tie-

molecules reside in the amorphous part. Because of the high molecular mass a large number of tie-molecules

is generated and the relatively long C2H5 side groups prevent extensive slipping of entanglement chains.

These two factors make the bonds between adjacent crystallites very stable. The crystalline lamellae form

spherical superstructures, consequently, a strong 3-dimensional network of entangled crystallites is formed

[34].

2.2.1 Mechanical properties

The Tab. 1 shows the mechanical properties of PB-1. The values shown are typical mid-range figures

and should not be considered as a specification – in fact, several different grades of PB-1 are produced with

properties tailored to satisfy different application requirements [47].

Tab. 1 Mechanical properties of PB-1 [34]

Material Properties Method Unit PB-1 4237

Tensile strength at yield ISO R 527 MPa 20

Tensile strength at break ISO R 527 MPa 35

Elongation at break ISO R 527 % 300

Flexural Elastic Modulus ISO 178 MPa 450

Notched Impact Strength at 20 °C ISO 180 kJ.m-2 20

Notched Impact Strength at 0 °C ISO 180 kJ.m-2 7

The peculiar tensile behavior of PB-1 is based mainly on the chain entanglements mentioned above.

PB-1 does not show the typical necking behavior; instead, it tends to support the load while it continues to

stretch. This is sometimes referred to as “ductile with work-hardening”. Depending on the preparation of the

test specimen and the conditions of measurement, PB-1 may exhibit a very little yielding in stress-strain

behavior (see Fig. 4).

24

Fig. 4 Tensile behavior of PB-1 vs. other polyolefins [34]

PB-1 has excellent impact toughness. The IZOD notched impact strength (ISO 180) of PB-1 is

classified “no break” at room temperature. The material retains its flexibility even at temperatures below

freezing point, which strongly facilitates installation during cold seasons. The glass transition temperature

determining the ductile/brittle transition of PB-1 lies at approx. –16 °C [34].

2.2.2 Physical properties

Density and melting point of various crystallographic forms of PB-1 shows Tab. 2.

Tab. 2 Density and melting point of PB-1

Material density g.cm-3

Tm oC ref

PB-1 amorphous 0.868 [58]

PB-1 form I 0.95 ~ 0.96

0.95 126 ~ 139

125

[59]

[58]

PB-1 form II 0.819 ~ 0.902

0.907

120 ~ 130

[59]

[58]

PB-1 form III 0.897 ~ 0.906 106 ~ 110 [59]

other polyolefins

PB-1

tens

ile s

tress

tensile strain

25

Isotactic poly(1-butene) has an excellent wet abrasion resistance, when tested in sand/slurry type

conditions. It performs as well as UHMW-PE which is well known for its outstanding abrasion and wear

resistance [34].

Thermal degradation of PB-1 results in decrease of molecular weight, melting temperature and glass

transition temperature [60].

2.2.3 Processing and application

However, even after more than 50 years from its discovery, and despite its potentials, PB-1 is still a

relatively new resin (particularly in respect to its application range) that is being produced in limited

quantities and has not yet found a large commercial success as compared to the other polyolefins. This is

likely due to both the relatively high monomer cost, possibly coupled with difficulties in polymerization

technology, and the peculiar crystallization behavior of the polymer [26].

Several possibilities of the polymer processing are reported [10]. PB-1 can be injection-molded,

blow molded, cast or blown into sheets and films, and extruded for wires coating [61]. The processing

behavior of PB-1 is somewhat intermediate between the behavior of HDPE and PP. Processing temperatures

are in the range 160 ~ 240 °C. Both die swell and cooling shrinkage are greater than for PE. The crystalline

material formed initially on cooling from melt is rather weak and must be handled with care on the haul of

equipment. As mentioned above the polymer must be aged for about a week in order to allow the more stable

crystalline form to develop [1].

Today, PB-1 is being discovered as a key to unlock opportunities in applications such as packaging,

construction, fibers and fabrics, compounds, adhesives, and coatings, just to name a few [3].

The main interest in PB-1 is in its use as a piping material, where the ability to use a lower wall

thickness for a given pressure requirement than necessary with other polyolefins, together with low density,

can lead in some cases to economic use. The principal application is for small-bore cold and hot water piping

(up to 95 °C) for domestic plumbing [1]. In Europe and Asia, PB-1 has been developed over the past 30

years to be arguably the technically preferred material. Pipe system manufacturers state that piping systems

made from PB-1 offer easy, fast installation with a reduced number of joints and connectors compared to

much stiffer conventional plumbing materials. In service, they note that the piping does not corrode or fur up

in hard water areas, it does not split at sub-zero temperatures and is quieter in use as it does not suffer from

"water hammer". It is used both in the construction of new buildings and the renovation of older properties.

A major UK housebuilder uses exclusively PB-1 piping systems. LyondellBasell's PB-1 is not being sold for

pipe applications intended for use in North America [2].

26

Further important application area for PB-1 is seal-peel or easy-open packaging. Typical examples

include carton liners (e.g. cereal packaging) and packs for pre-packed delicatessen products like cold meats,

cheeses and smoked salmon. PB-1 offers the ability to customize sealing temperatures and seal strength

whilst giving consistent, reliable processing and sealing performance on existing equipment, and built-in

tamper-evidence. PB-1 can also be used in film modification to increase flexibility and softness without

sacrificing clarity. In a similar vein, PB-1 is used to modify polypropylene fibers to enhance softness,

flexibility and to provide a unique feel [2].

PB-1 is also used commercially in films that require creep resistance such as insulation compression

packaging. It can also be used as a layer or overall film for hot filled heavy duty packaging, tapes for

bundling and holding loads at elevated temperature such as bale wrap, specialty sheet applications such as

anti-erosion geo-grids and netting and cable ties [2].

PB-1 features in hot melt adhesives, where its unique crystallization behavior reduces the melting

point and extends the "open time" of the adhesive to as much as 30 minutes, to enable time for repositioning

of components prior to setting. It is widely used as a component in hot melt adhesives for non-woven fiber

webs [2].

PB-1 is also typically used in applications that see continuous load at moderate to high temperatures.

It is used in electric domestic heaters in North America, where it is blow molded into large 15 ~ 20 kg tanks.

Other logical extensions may include fire extinguishers, residential and swimming pool water filter housings,

pressurized pneumatic holding tanks, hoses, compressed gas cylinders and aerosol dispensers, to name just a

few. There is also considerable interest in the potential of PB-1 to be used as a component in synthetic wine

corks [2].

PB-1 is also employed in compounded polymer products to provide unique benefits. It accepts

extremely high filler loadings (>70 %), which combined with its low melting point, enables it to be used in

halogen-free flame retardant composites or as a masterbatch carrier for difficult-to-disperse or temperature-

sensitive pigments [2].

27

2.3 Photodegradation Photodegradation (chain scission and/or crosslinking) occurs by the activation of the polymer

macromolecule provided by absorption of a photon of light by the polymer. In the case of photoinitiated

degradation light is absorbed by photoinitiators, which are photocleaved into free radicals, which further

initiate degradation (in non-photochemical processes) of the polymer macromolecules. In photo-thermal

degradation both photodegradation and thermal degradation processes occur simultaneously and one of these

can accelerate another. Photoageing is usually initiated by solar UV radiation, air and pollutants, whereas

water, organic solvents, temperature and mechanical stresses enhance these processes [62].

In spite of extensive studies of photodegradation of polymers, there are many unsolved problems.

Many publishers results are even contradictory and it is difficult to compare them, as the results are

frequently obtained on different samples of the same polymer and under different experimental conditions.

Many proposed mechanisms are still not experimentally proven [62].

Neat PB-1 because of contains only C-C and C-H bonds and is not, therefore, expected to absorb

light at wavelengths longer than 200 nm. The fact that photodegradation of polymers occurs even with light

>300 nm indicates that some kind of chromophoric groups must be present in these polymers. The

chromophores in commercial PB-1 are developed during polymerization and thermal processing.

During polymer photodegradation two types of processes occur: (i) primary photochemical reactions

due to the absorbed radiation, resulting in the formation of free radicals or non-radicals rearrangement; (ii)

secondary reactions in which radicals formed initiate a number of reactions which are independent of the

light [62].

As generally accepted, degradation reactions of semicrystalline polymers proceed predominantly in

amorphous regions. Nevertheless, physical factors, such as the size, arrangement and distribution of

crystalline regions, affect the degradation process as well. Photodegradation kinetics in polymer systems

depends on oxygen permeability through the material [63]. The rate of oxidation drops with decreasing

oxygen diffusion, following the increase of crystallinity and molecular orientation [64,65].

On the other hand, the crystallinity and the molecular orientation determine the mobility of the

radicals and therefore reduce the rate of termination, allowing an increase in the propagation of chemical

reactions leading to molecule scission, and this effect is opposite to that caused by reduced oxygen mobility.

Which of these two effects will predominate depends on the oxidation conditions, including the type of

degradation, temperature and oxygen concentration [65].

28

2.3.1 Radical oxidation

Hydroperoxides are well known to be crucial products of the free radical oxidation of most polymers

[66]. The mechanism of radical oxidation in poly(1-butene) is still unknown, nevertheless, the similar

mechanism as in the case of polypropylene is suggested, where resulting hydroperoxides are predominantly

tertiary (~ 90 %), the remainder being essentially secondary ones [67], on the contrary in polyethylene only

secondary hydroperoxides are present [68]. The general mechanism is shown in Fig. 5. These

hydroperoxides groups are the key products for understanding the mechanism as well as for gaining better

insight into the function of stabilizers [69]. Ouldmetidji et al. [70] found that peroxides in macromolecular

media can be detected using DSC which is more precise than classical titration of peroxides. Unfortunately,

in case of PB-1, its melting temperature is close to peroxide decomposition temperature, thus a modulated

DSC is required to detect the peroxide groups in PB-1.

initiation PH   ∆,   h     

 P·

chain propagation P· +O2

POO·

POO· + PH +O2

POOH + P·

chain branching

POOH     ∆,   h     

  PO· + ·OH

2 POOH    ∆,   h     

  PO· +H2O POO·

PO·

termination

P· + P·

   P· + POO·

POO· + POO·

+ O2

Fig. 5 Scheme of standard mechanism of the oxidation in polymer chain [62]

2.3.2 Aspects of natural weathering

Natural weathering offers reliable data on durability of stabilized polymers in a particular area of

application [71]. Nevertheless, this technique has some disadvantages. It is difficult to compare results

obtained in different geographical regions even when the materials are stressed by similar intensity of solar

radiation, however at different temperatures and atmospheric pollution [72].

Indeed, the weather as a whole is unpredictable and itself not a reproducible factor in short-term as

well as long-term periods. This is an unfavorable basis for attempts to predict long-term lifetime of polymers

affected by weathering. Basic data on weather components reflected in polymer weathering are briefly

mentioned here: solar radiation, oxygen and casual oxidizing and acid atmospheric pollutants are principal

atmospheric detergents responsible together with environmental temperature and humidity for the reduced

durability of outdoor exposed stabilized polymers. Microorganisms add to the failure. Various detergents act

in concerted processes and their effects are mostly complementary or even synergistic [73].

29

UV radiation of 295 ~ 400 nm characteristic of a continuous spectrum and constituting about 1 ~ 5 %

of the total solar radiation is a dangerous component initiating most outdoor processes by photolysis of

polymer or additive covalent bonds, photolysis of hydroperoxides and excitation of chromophores [62,74]. It

follows that an increase in intensity of UV radiation or involvement of more energetic UV-B radiation below

295 nm in artificial radiation sources accelerates photoprocesses, but may at the same time cause differences

in degradation mechanisms and stabilizer durability. Visible light (400 ~ 760 nm) forms 39 ~ 53 % of the

terrestrial solar radiation and increases photodegradation by sensitization of visible light-absorbing

chromophores or additives. Besides UV radiation, some level of visible light emitted by artificial sources in

accelerated tests is recommended. Infrared radiation (760 ~ 2500 nm, 42 ~ 60 % of the solar radiation)

triggers thermal effects, particularly in dark pigmented polymers [74]. Infrared radiation is a source of

differences measured between environmental temperature and surface temperature of the irradiated polymer.

Formally, an exact simulation of the ‘‘standard’’ solar radiation distribution in well-designed accelerated

tests is possible. However, under natural conditions, sunlight emissions vary in spectral energy distribution

and intensity [71]. Even on using optimized artificial radiation sources (e.g., filtered xenon arcs) an exact

simulation of natural exposure is hindered. Nevertheless, application of devices with artificial sources makes

possible a comparison of results between various laboratories [73].

Besides oxygen, the common oxidizing component, the Earth’s atmosphere contains traces of

pollutants arising from biogenic or anthropogenic emissions. Ozone and nitrogen oxides (NOx) are oxidizing

and reducing pollutants co-acting with solar radiation. Their atmospheric photochemistry as well as details of

their effects on rubbers is well understood [75,76]. Interactions with plastics or stabilizers in particular are

not fully explained. They can be generally categorized as phototriggering impurities [74]. Moreover, NOx are

responsible for photo-assisted depletion of phenolic antioxidants (in gas fading) [77].

The effect of pressurized chlorinated water in PB-1 pipes was studied by Lundback et al. [78]. It was

observed that antioxidant concentration is independent of chlorine concentration and the lifetime shortening

was approx. by a factor 10 with respect to that obtained in pure water.

PB-1 is sensitive to oxidizing acids, aromatic and chlorinated hydrocarbons. The polymer is very

sensitive to environmental stress cracking [2].

2.3.3 Crystallization of degraded polyolefins

The crystallization of polymers depends on many factors including the shape and stereoregularity of

the repeat unit, the presence of branches and crosslinks, and the molecular chain length in the linear

segments. In polypropylene, the repeat unit is small and can easily be assembled into crystal form. When

polypropylene is subjected to photo-oxidation several changes occur that influence the crystallinity and the

crystallizability of the material [65,79]. Caig et al. [79] studied oxidation that occurs at or near to ambient

30

temperature, when the polymer is in solid form, well below the crystallization temperature. The changes

occur predominantly in the non-crystalline phase because oxygen can diffuse through such regions relatively

freely but is almost excluded from the crystalline regions. The principal changes are (i) chain scission; (ii)

crosslinking; and (iii) the formation of molecular defects such as carbonyl groups. Chain scission releases

chain segments that were previously entangled and allows them to crystallize. If enough material of this kind

becomes available to crystallize, new crystals could form but it is more likely that these segments will attach

to the growth faces of pre-existing crystals nearby. This will increase crystallinity and is a form of secondary

crystallization often known as “chemi-crystallization” [64,80,81]. Crosslinks inhibit further crystallization of

the chain segments that are connected by them. Molecular defects such as carbonyl groups, which are

formed, do not fit into the crystal lattice and the parts of the molecular segments containing them will not be

able to take part in secondary crystallization [79].

Chemi-crystallization is caused by a variety of molecular degradation processes and has been

detected in many polymers [82-86]. If, as expected, oxidation within the crystal phase is relatively rare, the

network effect of photo-oxidation will be an increase in crystallinity and the formation of material in the

amorphous phase that becomes increasingly unable to crystallize even if scission events or elevated

temperature mobilize the chain segments. Thus, it is expected that, after a certain exposure, the rate of

chemi-crystallization will decay to zero. Chemi-crystallization is an important phenomenon because the

density of the crystals is higher than that of the amorphous material and shrinkage occurs. Under UV

irradiation, the scission (and hence secondary crystallization and, in turn, shrinkage) tends to vary very

sharply with distance from the exposed surface [87,88] and this leads to the development of tensile residual

stresses near the surface [89,90]. The residual stresses and also freezed stresses arising from a processing

[91] can lead to cracking of the material that has already been embrittled by the molecular degradation,

through the loss of entanglements and the consequent loss of mechanical integrity provided by the molecular

network, and through the reduced flexibility that results from increased crystallinity and crosslinking that

may be present additionally [87,88]. Serious deterioration of engineering properties may therefore occur. If

the cracks are not very deep and do not lead to total failure, the appearance may become unacceptable

because of fine surface cracks that spoil the glossy appearance, etc. Residual stresses lead to distortion if they

are not balanced across the molding section.

When photo-oxidized material is melted and cooled down again, the crystallinity of the newly

solidified material will depend not only on the polymer and the cooling conditions but also on the molecular

changes that occurred during photo-oxidation. The shorter chains produced by scission events (which are still

regular) will crystallize more readily whereas crosslinks and molecular defects will not be able to crystallize

and will be rejected from the newly formed crystals. Thus, there are two opposing effects, one that promotes

greater crystallizability and the other inhibiting crystallization [79].

31

3 SUMMARY AND AIMS Isotactic polybutene-1 is a polymorphic material similar to isotactic polypropylene, however its

peculiar transformation behavior is a main factor inhibiting its use as compared the other polyolefins, namely

polyethylene and polypropylene. As written above in Chapter 2, PB-1 crystallizes from melt to kinetically

favored form II and according to annealing conditions it recrystallizes to thermodynamically stable form I.

The transformation behavior of technologically processed materials has not practically been studied yet and

the degradation behavior of PB-1 is unknown as well. The lack of knowledge can be considered as one of the

reasons explaining rather low practical use of PB-1.

Therefore, this thesis is concentrated on study of transformation behavior in processed PB-1 under

various thermodynamic conditions and changes in structure and properties are observed; the degradation

behavior of naturally weathered and accelerated UV irradiation.

The main aims of this thesis are according to present knowledge following:

1. Phase transformation in PB-1 upon various thermodynamic conditions.

2. Evolution of structure and properties of injection-molded and extruded PB-1.

3. The effect of photodegradation on the evolution of structure and properties in PB-1.

32

4 MATERIALS AND EXPERIMENTAL TECHNIQUES

4.1 Materials In this work, three commercially-available grades of isotactic poly(1-butene) produced by Basell

Polyolefins, Louvain la Neuve, Belgium were used. The basic properties of PB-1 grades are described in

table Tab. 3.

Tab. 3 Used materials of PB-1 [92]

Material grade – method PB 8640M DP 0401M PB 0300M

Characteristics Random copolymer of butane-1 and ethylene

Semi-crystalline homopolymer

Semi-crystalline homopolymer

Density (g.cm-3) - ISO 1183 0.906 0.915 0.915

Melt flow index - ISO 1133 (190 °C / 2.16 kg)

1 g/10 min 15 g/10 min 4 g/10 min

Flexural modulus - ISO 178 250 MPa 450 MPa 450 MPa

Tensile strength at yield ISO 8986-2

22 MPa 19.5 MPa

Tensile strength at break ISO 8986-2

30 MPa 29 MPa 35 MPa

Tensile elongation at break ISO 8986-2

300 % 300 % 300 %

Melting temperature m1 DSC

113 °C 126 °C 127 °C

Melting temperature m2 DSC

97 °C 114 °C 116 °C

Mechanical properties were measured on specimens conditioned for 10 days at 20 °C

4.2 Sample preparation

4.2.1 Compression molding

Sheets with thickness of 0.3 and 1.0 mm were prepared by compression molding at temperature of

190 °C, compression time of 5 minutes and subsequent cooling to 20 °C for 5 minutes.

4.2.2 Injection molding

A DEMAG NC 4 injection molding machine was employed for production of tensile testing

specimens according to standard ISO 527 (tensile dumbbell test bars with a gauge length of 80 mm, width of

10 mm and thickness of 4 mm). The main processing parameters were following for all material grades:

Temperatures in the barrel: 50, 170, 190, 200 and 210 °C (feed, transition, metering and nozzle); injection

pressure 75 MPa with injection speed 50 mm/s; holding pressure 50 MPa with time 25 s; cooling time 40 s;

mold temperature 50 °C.

33

4.2.3 Extrusion

The materials were extruded using a Brabender extruder with temperature setting 140 °C, 145 °C,

150 °C (barrel) and 150 °C (head) . A rotation speed was 20 rpm. The extruded tape had approx. 20 mm

width and thickness approx. 2 mm. Then, the tapes were cut off to the pieces with length 160 mm.

4.3 Ageing

4.3.1 Annealing

The prepared specimens were immediately after processing annealed at various temperatures: -22,

+5, +22, +40 and +60 °C at atmospheric pressure for various times. The temperatures were chosen with

respect to utilization of PB-1, which can be used for tanks, hose, tubing, molded parts, films, etc. at various

temperatures. The chosen temperature of -22 °C is generally temperature in freezer; +5 °C refrigerator, cold

water; +22 °C room temperature; +40 °C hot weather temperature, water; +60 °C hot water, heating.

4.3.2 Natural weathering

Natural weathering was performed in Nivnice, Czech Republic, altitude of 250 meters, geographical

position 48°58′42″N and 17°38′47″E. Specimens were exposed in south direction under the exposition angle

of 45°, between May 7 and July 13, 2008, i.e. 0 ~ 63 days in case of sheets 0.3 mm thick and during June 6

to September 17 2008 i.e. 0 ~ 103 days in case of dog-bone specimens.

The meteorological data are in the table 4.

Tab. 4 meteorological data [93]

Month rain-fall

(mm) T (°C) clearness

day

cloudiness

day

thunderstorm summer day tropic day

May 82.2 15.2 9 4 6 12 1

June 78.8 19.8 10 1 17 14 10

July 104.7 19.7 10 14 8 12 10

August 95.3 18.9 14 1 12 18 7

September 62.3 13.7 9 9 1 9 3

Explanation: T - average temperature (calculated by three temperatures at 7, 14 and 21 hours; evening observation is calculated twice and divided by 4); clearness day – max 20 % of covered sky; cloudiness day – cloudy 80 – 100 %; summer day – part of day >25 °C; tropic day – part of day >30 °C

4.3.3 UV irradiation

For accelerated UV irradiation, a SEPAP 12/24 MPC type of commercial device was employed at

temperature of 60 °C, UV irradiation wavelength λ > 300 nm and the sampling was performed in various

ageing times.

34

Each quarter angle of rotating carousel with specimens is covered by one lamp vapor mercury

“medium pressure” (MAZDA 400W) (Fig. 6), en emission of spectra of which is polychromatic. The

wavelength shorter than 300 nm is filtered by Pyrex glass cover adjustment of lamps where its emission of

spectra is superposed to the sun spectra (Fig. 7). The carousel with specimens rotates with speed 4 rotations

per minute to assure the homogenization of the luminous flux by specimens. Distance between the specimen

and the lamp is 20 cm. The temperature on the surface of the specimens is controlled by a platinum sonde

(Pt 100) which is in contact with PE film fixed on carousel. Temperature regulation is assured by a

ventilator; the temperature was fixed to 60 °C.

Fig. 6 Scheme of a SEPAP 12-24

Fig. 7 Spectra of lamps vapor Hg medium pressure used in a Sepap 12-24

 

200 300 400 500 600 700 800

0

2

4

6

8

10

12

14

16

18

inte

nsity

(u.a

.)

wavelength (nm)

Hg Hg

Hg Hg

carousel holder of specimens

20 cm

35

4.4 Testing and characterization

4.4.1 Tensile testing

Tensile testing was performed using an INSTRON 8871 apparatus at room temperature. The

specimens with a gauge length of 80 mm were deformed to a maximum elongation Δl = 100 mm. The rate of

extension was 100 mm/min.; tensile modulus, yield stress and extension at break were calculated by a

Bluehill software, which is a part of the tensile testing apparatus for its operating and evaluation of the

results. Five specimens were tested and average values are reported.

In addition, a Zwick 145665 multipurpose tester was used for the tensile testing of extruded tapes

with a gauge length of 80 mm (specimen dimensions - width x thickness x length – 20 x 2 x 110 mm) and

dog bone-shaped specimens (ISO 527-3 type 5A) cut out from the tapes with a gauge length of 30 mm. The

specimens were strained at room temperature up to break at a test speed of 50 and 100 mm/min. From the

stress–strain traces, yield stress, yield strain, stress at break and strain at break were derived. Besides, the

elastic modulus was evaluated using a Zwick external extensometer (gauge length of 20mm) at a test speed

of 1mm/min. Five specimens were tested and average values are reported.

4.4.2 Polarized light microscopy

The spherulitic morphology of the specimens was observed using a Zeiss NU light microscope

equipped with a polarizer. For this purpose, thin slices with thickness approximately of 40 μm were cut out

using a Leica rotational microtome. The observed morphologies were taken by Sony DSC F717 camera.

4.4.3 Stereomicroscopy

The surfaces of the degraded materials were surveyed by Stereo microscope ZEISS Stemi 2000-C

equipped by Sony DSC F717 camera.

4.4.4 Scanning electron microscopy

Scanning electron microscopy (SEM) is useful method for most types of surface morphology

examination. In this work, it was used for analyzing of degraded surfaces.

The surface of the specimens were coated with a thin gold/palladium (SC7620 Mini Sputter Coater)

film and examined in a scanning electron microscope Vega\\LMU (Tescan, Czech Republic).

4.4.5 Wide-angle X-ray scattering

Wide-angle X-ray scattering (WAXS) is the traditional method of crystallographic structure

determination, the standard technique and analysis procedure, which can be used in the study of crystalline

polymers. WAXS measurements were used to determine the degree of crystallinity and polymorphic

composition of the prepared specimens.

36

For scattering in the transmission mode, diffraction angle with interval 2θ = 7 ~ 25°, the measuring

step of 0.05° and holding time of 5 s were used. A URD6 diffractometer CuKα radiation monochromatized

with a Ni filter (λ = 0.154 nm) was employed for measuring of PB-1 extrudates in reflection mode. Whole

extrudates (only cut in length) were used as specimens for scattering.

For scattering in the reflection mode, a Philips Xpert was employed with interval 2θ = 5 ~ 30°,

measuring step of 0.05° and holding time of 0.5 s, CuKα radiation monochromatized with a Ni filter (λ =

0.154 nm). The degraded sheets with thickness of 1 mm and the degraded extrudates were scattered by this

device.

The reflections relative to the tetragonal form of PB-1 are basically three main peaks at 11.9, 16.9

and 18.48 2θ, corresponding to the (200), (220) and (301) planes. The hexagonal phase I is characterized by

four diffraction peaks at 9.9, 17.3, 20.2 and 20.58 2θ, originated by the (110), (300), (220) and (211) planes,

respectively [20]. It may be noted that, as a consequence of the phase transition, with increasing time the

intensity of the peaks typical of the hexagonal phase I (i.e.(110) at 9.98 2θ) increases, whereas the intensity

of peaks related to the tetragonal phase II (i.e. (200) at 11.98 2θ) decreases (see Fig. 8).

A PEAKFIT v4 was used for the evaluation of crystalline part in the specimens. It is suggested to

follow the disappearing of the peak at 11.8° 2θ (reflection plane 200) the form II rather than increase the

peak at 9.9° 2θ (reflection plane 110) of the form I because, together with the phase transformation, a post

crystallization phenomenon takes place during which part of the amorphous polymer crystallizes into the

form I [26]. However, in this work, it was not possible to measure the identical specimen but a series of

specimens prepared under controlled conditions. It was decided to follow evolution of the both peaks of

forms I and II (at 9.9° and 10.8° 2θ) and calculate the ratio of heights both peaks and content of the form I

from the total sum of both heights. The same approach has been used in several works by Azzuri et al. [51],

Samon et al. [94] and Natta et al. [8]. The similar approach was used in work of Kaszonyiova et al. [37] or in

case of isotactic polypropylene, which is also a polymorphic material [95,96].

Relative crystallinity cx was calculated according to the equation (1) as a ratio of integral intensity

diffracted in certain measuring direction by crystalline part Ic and a total intensity I:

   ⁄ · 100 % ( 1 ) 

4.4.6 Me

It is

parameters

changes of

(MWD). Th

which invol

Bas

Mo

using rheolo

time depend

No

The comple

elt rheology

s well know

through stru

f such mole

hrough phot

lve cross-link

sic melt visco

st polymeric

ogical exper

dence of mec

single param

ex dynamic m

Fig. 8 Wide-

wn that melt

ucture relation

cular param

to-ageing, th

king in case

oelasticity [9

c materials e

iments such

chanical prop

meter can be

modulus (G*

-angle X-ray s

viscoelastic

nships. Henc

meters as mo

hese changes

of elastomer

97]:

exhibit a vis

as dynamic

perties of pol

used to char

) is resolved

scattering patt

ity measurem

ce, the evolut

olecular wei

s can be cor

rs [76].

scoelastic be

mechanical

lymers.

racterize the

d into two com

i "

terns at differe

ments can p

tion of rheol

ight (MW),

rrelated to ch

ehavior. Visc

testing, whic

stress-strain

mponents us

ent times [26]

provide infor

logical prope

and molecu

hain scission

coelastic pro

ch offers a c

relationship

sing complex

rmation abou

erties directly

ular weight

n and/or rec

operties are

convenient w

p in viscoelas

x notation:

37

ut molecular

y reflects the

distribution

ombinations

investigated

way to assess

stic material.

( 2 )

7

r

e

n

s

d

s

.

)

38

The real part of the complex modulus (G’) describes stress-strain relationships that are in phase. G’

is called the storage modulus (or elastic modulus). The imaginary component (G”) characterizes the out-of-

phase component and is named the loss modulus (or viscous modulus). Dynamic viscosity (η*) is related to

the complex modulus by

d d⁄⁄ i i "⁄ ( 3 )

with "/ and ” / . Then, the real component of the complex viscosity (η’) describes

the viscous dissipation in the specimen, while the imaginary component (η”) represents the stored elastic

energy. Furthermore, the tangent of the phase angle (tan δ) describes the balance between the viscous and

elastic behaviors in a polymer melt:

  tan "⁄ "⁄ ( 4 )

It is well known that the evolution of the rheological material properties directly reflects changes in

molecular parameters. The linear viscoelastic properties in dynamic experiments are sensitive both to the

chain scission and to the three-dimensional network formation. Thus, melt rheology provides a convenient

tool to view the particular behavior due to the competition of chain scissions and recombination reactions

occurring through ageing.

It is also well known that the zero shear viscosity η0 depends on the molecular weight and obeys a

power law [98,99]:

  0      ( 5 ) 

with the widely quoted viscosity exponent α = 3.4.

The zero shear viscosity η0 can be obtained from the complex viscosity η*(ω):

  ω i i "⁄   ( 6 ) 

and

  | |0 0 0 ( 7 ) 

An empirical rheological model used to fit dynamic data is the Cole-Cole distribution expressed by

[100-104]:

  01 i 0 1 ( 8 ) 

where λ0 is the average relaxation time and h is a parameter of the relaxation-time distribution.

In the complex plane this model predicts the variation of the viscosity components (η’ versus η”) to

be an arc of circle. From this representation it is easy to determine the parameters of the distribution: η0 is

39

obtained through the extrapolation of the arc of the circle on the real axis and the distribution parameter h

through the measurement of the angle Φ Π/2 between the real axis and the radius going from the origin

of the axis to the centre of the arc of the circle.

Molecular changes were characterized by melt viscoelasticity experiments in oscillatory shear mode

using a TA ARES mechanical spectrometer, equipped with parallel plate geometry with diameter of 8 mm

and 25 mm and the gap between the plates was 1 mm. The experiments were carried out at temperatures of

120, 140 and 160 °C.

IRIS 2006 (Innovative Rheological Interface Software) was used to create master curves by the

automatic time-temperature superposition at temperature of 140 °C.

4.4.7 Infrared spectroscopy

For the purpose of identification the chemical changes of PB-1 through UV irradiation, several

devices were employed.

A Nicolet Impact 760 spectrometer in FTIR transmission mode, with nominal resolution of 4 cm-1

and 32 acquisitions was employed.

A Nicolet 380 FT-IR spectrometer equipped with SPECAC ATR Diamand-Germanium crystal, with

nominal resolution of 4 cm-1 and 32 acquisitions, was employed for surface analysis of degraded specimens.

Omnic software was used for collection of all spectra.

4.4.8 Differential scanning calorimetry

Melting, remelting and crystallization behaviors of UV irradiated specimens were studied using a

power-compensated differential scanning calorimeter (DSC - Pyris1, The Perkin-Elmer Co., USA). Nitrogen

as a purge gas, was used and constantly passed (20 ml/s) through the heat sink and over the cells.

Temperature calibration was performed using indium as a standard. Approx. 5 mg of irradiated material was

loaded into standard aluminum pans. Observations were performed at a rate of ±10 °C/min in interval from

40 to 170 °C. To observe melting temperature of the form II, which is present after melting and transforms in

time into form I, second subsequent melting and crystallization was performed. The specimens were kept at

room temperature at 23 °C and the phase transformation followed.

It is generally known that differences in melting temperatures of the forms I and II are about of 10 ~

15 °C. Thus melting curves obtained from DSC exhibit partial superposition. Only one melting peak of the

form I or II on the melting curve is detected immediately after crystallization or when polymorphic

transformation terminates respectively.

In t

Therefore, w

It i

Unfortunate

Thu

software us

was found t

contents of

of both pe

transformat

this respect,

we particular

s supposed t

ely, there are

us, we integr

sing a Gauss

that ratios of

the increasin

eaks. Also t

ion is compl

Fig

the melting

rly followed

that overall d

e no exact va

rated the are

sian function

f the deconv

ng form I in t

the ratios of

leted the ratio

g. 9 Deconvol

g curves mu

the approach

degree of cry

lues of melti

as of the pe

n. Then, area

voluted peak

time were ca

f the forms

os are zero o

lution of the D

ust be decon

h of Alfonso

ystallinity do

ing enthalpy

aks of the fo

as and heigh

ks areas are e

alculated as t

I and II w

or infinitive.

DSC curve obt

voluted to r

o et al. [105].

es not chang

of forms I an

orms I and I

hts of individ

equaled to th

the ratio heig

were calcula

tained by PEA

reliably evalu

ge during tran

nd II.

I deconvolut

dual melting

he heights of

ght of peak fo

ated, howev

AK-FIT V4 [1

uate the tran

nsformation

ted via a PE

peaks were

f the melting

orm I and sum

er in time

05]

40

nsformation.

[38,43,105].

EAK-FIT V4

obtained. It

g peaks. The

m of heights

0 and after

0

.

.

4

t

e

s

r

41

5 STUDY OF TRANSFORMATION IN PB-1 Due to of manufacturing PB-1 parts from its melt, the resulting properties are affected by the phase

transformation. Therefore, the understanding of the transformation is desirable in connection with polymer

applications. The purpose of this study was to characterize influence of the temperature on phase

transformation and evolution of mechanical properties in injection-molded and extruded PB-1 specimens.

5.1 DSC study The phase transformation behavior was in the first step observed by DSC experiments.

The crystallization temperatures of the used materials, done on compression-molded specimens, during non-

isothermal crystallization at cooling rate of 10 °C/min are following: DP 0401M = 78 °C; PB 0300M =

84 °C; PB 8640M = 51 °C.

In the following Figs 10-12 the temperature dependent evolution of form I content in all three

materials can be seen. The evolutions in selected days for all three materials in whole temperature range are

illustrated in Fig. 13.

What can be seen in all materials is that fastest phase transformation occurs at temperatures of +5

and +22 °C. At temperature of +40 °C, a slower transformation is observed in PB 0300M and PB 8640.

It was already found that addition of ethylene copolymer to PB-1 increases phase transformation rate; this

fact is confirmed in case of ethylene copolymer PB 8640M at temperature of -22 °C. This can be actually

related to a decrease of glass transition temperature by the ethylene addition which can still keep the

molecular mobility in the system. The reason, why the fastest transformation rate is not observed at higher

temperatures (+40 and +60 °C) can be explained by an influence of annealing at temperature above the

crystallization temperature, which is in case of PB 8640M 51 °C, while in case of homopolymers (PB 0300M

and DP 0401M) it lies above annealing temperatures. Probably post-crystallization occurs at this temperature

of 60 °C in PB 8640M, which can cause a competition the original crystallite perfection and the phase

transformation and thus, it is slowed down.

In case of both homopolymers at temperatures of -22 and 60 °C similar behavior can be seen with

the difference in the reached amount of the form I content which is in PB 0300M higher approx. about 10 %

compared to DP 0401M. This can be caused by different MW, which is higher in PB 0300M than in

DP 0401M.

42

Fig. 10 Evolution of form I content in DP 0401M for various annealing temperatures

Fig. 11 Evolution of form I content in PB 0300M for various annealing temperatures

0 2 4 6 8 10 12 14 16 26 27 28

0

10

20

30

40

50

60

70

80

90

100

cont

ent o

f for

m I

(%)

annealing time (days)

-22 °C +5 °C +22 °C +40 °C +60 °C

0 2 4 6 8 10 12 14 16 26 27 28

0

10

20

30

40

50

60

70

80

90

100

cont

ent o

f for

m I

(%)

annealing time (days)

-22 °C +5 °C +22 °C +40 °C +60 °C

43

Fig. 12 Evolution of form I content in PB 8640M for various annealing temperatures

0 2 4 6 8 10 12 14 16 26 27 28

0

10

20

30

40

50

60

70

80

90

100

cont

ent o

f for

m I

annealing time (days)

-22 °C +5 °C +22 °C +40 °C +60 °C

44

Fig. 13 Evolution of form I content in the whole temperature range for all materials in selected annealing days

-20 -10 0 10 20 30 40 50 600

10

20

30

40

50

60

70

80

90

100

-20 -10 0 10 20 30 40 50 600

10

20

30

40

50

60

70

80

90

100

-20 -10 0 10 20 30 40 50 600

10

20

30

40

50

60

70

80

90

100

-20 -10 0 10 20 30 40 50 600

10

20

30

40

50

60

70

80

90

100

cont

ent o

f for

m I

DP 0401M PB 0300M PB 8640M

day 1 day 2

cont

ent o

f for

m I

(%)

annealing temperature (°C)

day 4

annealing temperature (°C)

day 9

45

5.2 Evolution of properties of injection-molded PB-1

5.2.1 Tensile modulus

Results of tensile testing (elastic modulus) for all materials are illustrated in Figs 14-16. The non-

monotonic evolution of tensile modulus in dependence on annealing temperature is obvious.

In all materials, the fastest evolution of tensile modulus is evident at temperature of +5 and +22 °C,

followed by a bit slower evolution at temperature of +40 °C. This is in agreement with previous Chapter 5.1,

and below or above +5 ~ +22 °C it slows down. The slowest evolution was observed in the specimens

annealed at temperature of +60 °C and -22 °C with distinct differences.

The differences between homopolymers (Figs 14 and 15) and PB 8640M (Fig. 16) can be seen, even

between two homopolymers.

PB-1 ethylene copolymer PB 8640M exhibits faster evolution at temperatures -22 and +60 °C

compared to homopolymers, which is in agreement with previous findings. Also, as for the other

temperatures it shows notable higher increase in elastic moduli in the first two annealing days compared to

homopolymers, nevertheless the further evolution is slow and the maximum tensile modulus are reached

after 14 annealing days, particularly at temperature of +60 °C.

Apparent differences were observed in homopolymers as the transformation of PB 0300M is slower

as compared to DP0401M at temperatures of +5 and +22 °C, where virtually maximum moduli are observed

after 4 days of annealing, at temperature of +40 °C after 7 days. However, in PB 0300M the phase

transformation is more gradual in whole temperature range with the fastest transformation at +5 °C. The

material PB 0300M possesses higher molecular weight compared to DP0401M, which probably may play a

role in evolution. Indeed, its molecular weight and consequent number of molecular entanglements is higher

as compared do DP 0401M; this limits the system mobility and can results in gradual transformation rate. On

the contrary at temperatures of -22 and +60 °C, the phase transformation is notably faster in PB 0300M.

In both homopolymers the tensile modulus slightly increases even after 48 days, which is more

pronounced at temperature of -22 and +60 °C. This can be caused by a post-crystallization or by a slow

continuous transformation phase II to I.

To obtain detail insight into phase transformation the other tensile properties were also evaluated as

yield strength, elongation at break and strength at break.

46

Fig. 14 Evolution of tensile modulus in injection-molded PB 0300M

Fig. 15 Evolution of tensile modulus in injection-molded DP 0401M

0 5 10 15 20 25 45 46 47 480

20

40

60

80

100

120

140

160

180

200

220

tens

ile m

odul

us (M

Pa)

annealing time (days)

-22 °C +5 °C +22 °C +40 °C +60 °C

0 5 10 15 20 25 45 46 47 480

20

40

60

80

100

120

140

160

180

200

220

240

tens

ile m

odul

us (M

Pa)

annealing time (days)

-22 °C +5 °C +22 °C +40 °C +60 °C

47

Fig. 16 Evolution of tensile modulus in injection-molded PB 8640M

5.2.2 Tensile yield strength

The tensile yield behavior was evaluated only in DP 0401M (see Fig. 17) for temperatures of +5,

+22 and +40 °C since the annealing day 4 resp. 7. It can be seen that yield strength increases with increasing

annealing time similarly as the tensile modulus. It can be supposed that observation of yield behavior is

affected by a molecular weight; from all three materials DP 0401M possesses the lowest one. In addition, the

observation of yield strength can be probably affected by a critical amount of the form I, respectively by a

minimal modulus.

Fig. 17 Evolution of tensile strength at yield in injection-molded DP 0401M

0 5 10 15 20 25 45 46 47 48 490

20

40

60

80

100

120

140

160

180

200

220

240

tens

ile m

odul

us (M

Pa)

annealing time (days)

-22 °C +5 °C +22 °C +40 °C +60 °C

0 5 10 15 20 25 30 35 40 45 5017.0

17.5

18.0

18.5

19.0

19.5

20.0

20.5

21.0

yiel

d st

reng

th (M

Pa)

annealing time (days)

+5 °C +22 °C +40 °C

48

5.2.3 Tensile elongation at break

The tensile elongation at break was observed only in materials PB 0300M and PB 8640M (see Figs

18 and 19), because of the limitation of extension on tensile testing apparatus. It has in the whole

temperature range, except the temperature of -22 °C, decreasing trend during annealing. It is opposite effect

in respect to the evolution of tensile modulus.

Nevertheless, at annealing temperature of -22 °C a specific behavior is observed in both materials. It

can be supposed that evolution in PB 0300M possesses similar behavior as in PB 8640M, where an

unexpected increase of elongation at break occurred compared to annealing temperature of +60 °C. Between

the days 14 and 24, significant drop to a half value was observed in PB 8640M and further annealing lead to

virtually the same value as at temperatures +5, +22 and +40 °C. We suppose that longer annealing of PB

0300M would probably lead to equilibrating of elongation at break for the other annealing temperatures.

Fig. 18 Evolution of tensile elongation at break in injection-molded PB 0300M

0 2 4 6 8 10 12 14 16 18 20 22 24 46 48

50

60

70

80

90

100

110

120

130

tens

ile e

long

atio

n at

bre

ak (%

)

annealing time (days)

-22 °C +5 °C +22 °C +40 °C +60 °C

49

Fig. 19 Evolution of tensile elongation at break in injection-molded PB 8640M

0 5 10 15 20 25 46 48 50

30

40

50

60

70

80

90

100

110

120

130

tens

ile e

long

atio

n at

bre

ak (%

)

annealing time (days)

-22 °C +5 °C +22 °C +40 °C +60 °C

50

5.2.4 Tensile strength at break

Conversely, the evolution of tensile strength at break gradually rises with increasing annealing time

as can be seen in Figs 20 and 21. In both materials, annealing temperature of +5 °C lead to the highest values

of strength at break. When both Figs are compared with those of tensile modulus (Figs 14 and 16), similar

trend of the evolution can be seen. While in the case of tensile yield strength (in DP 0401M - Fig. 17) the

comparing with tensile modulus (Fig. 15) was not possible.

Fig. 20 Evolution of tensile strength at break in injection-molded PB 0300M

Fig. 21 Evolution of tensile strength at break in injection-molded PB 8640M

0 5 10 15 20 25 46 48 5013

14

15

16

17

18

19

20

21

22

tens

ile s

treng

th a

t bre

ak (M

Pa)

annealing time (days)

-22 °C +5 °C +22 °C +40 °C +60 °C

0 5 10 15 20 25 46 48 508

9

10

11

12

13

14

15

16

17

18

19

20

21

22

tens

ile s

treng

th a

t bre

ak (M

Pa)

annealing time (days)

-22 °C +5 °C +22 °C +40 °C +60 °C

51

5.3 Evolution of structure and properties in PB-1 extrudates Since PB-1 is also processed by extrusion (mostly for pipe applications), it opens an industrial

question whether different processing technology and subsequent morphology can play a role in phase

transformation. In comparing with previous Chapter 5.2 with injection-molded specimens, it has been

extented by WAXS data.

5.3.1 WAXS evolution

The wide-angle X-ray scattering was performed in transmission mode to observe the evolution of the

form I and II in PB 0300M as can be seen in following Fig. 22. It shows a virtually the same evolution of the

form I content at five different annealing temperatures. It can be seen that fastest evolution is observed at

annealing temperature of +5 °C and from day 14 the form I content becomes constant at approx. 90 %.

At temperature of +22 °C the evolution is a rather slower, however during first 3 days its rate is the same as

at +5 °C. Gradual evolution occurs at temperature of +40 °C, which is slower compared to +5 and +22 °C

but with a similar trend. On the contrary, a peculiar trend of the evolution occurred at annealing temperatures

of -22 and +60 °C. Until the day 14, a slow gradual evolution seems to be similar for both temperatures, with

respect to peculiar increase and decrease at -22 °C. Nevertheless further observation at time of 22 and 30

days shows unexpected significant increase of the form I content to over 80 %. The confirmation of this

behavior was found in subchapter 5.3.3 – Tensile modulus. This can signify that at annealing temperatures of

-22 °C an induction time was observed, during which slow transformation occurs until critical amount of the

form I is transformed. The reached critical amount can probably accelerate the transformation rate.

Nevertheless, this hypothesis cannot be applied for injection-molded specimens, where the phase

transformation was faster for +60 °C than for -22 °C. However, this probably signifies that processing

technology respectively resulting morphology directly affects the phase transformation during temperature

annealing.

52

Fig. 22 WAXS evolution of the form I content in extruded PB 0300M for various annealing temperatures upon time

0 5 10 15 20 25 30 55 56 57

0

10

20

30

40

50

60

70

80

90

100

cont

ent o

f for

m I

(%)

annealing time (days)

-22 °C +5 °C +22 °C +40 °C +60 °C

53

5.3.2 Polarized light microscopy

The morphology of extruded specimens (Fig. 23) shows differences between the materials; isotropic

spherulitic structure can be seen. It should be noted that extrusion conditions were the same for all materials.

The main difference can be seen in the morphology of homopolymers. Material DP 0401M possesses lower

MW and this can probably affects the morphology – remarkably smaller and finer spherulite structure when

compared to PB 0300M – the growth rate is apparently higher than the nucleation rate.

Fig. 23 Images of morphology in extruded PB-1(bottom is border of specimen)

54

5.3.3 Tensile modulus

The evolution of tensile modulus in all materials is illustrated in following Figs 24-26.

When previous WAXS observation (Fig. 22) and evolution of tensile modulus (Fig. 24) of material

PB 0300M are compared, distinct similarities can be seen: (i) a peculiar increase occuring at annealing

temperature of -22 °C is observed in both Figs; it crosses at the same time (approx. 28 days) the evolution

curve at annealing temperature of +40 °C; (ii) annealing temperature of +40 °C possesses virtually the same

trend of gradual evolution; (iii) low content of form I in specimens annealed at +60 °C follows its low tensile

modulus; (iiii) rather higher content of the form I in specimens annealed at +5 °C is directly reflected in

higher tensile modulus compared to specimens of +22 °C.

As suggested in previous subchapter 5.3.1 WAXS evolution - an induction transformation time

during which critical amount of form I is transformed, seems to be relevant. After reaching the critical

amount of the form I, a rapid increase is initiated and plato is reached according to annealing temperature.

A comparison of all the materials shows the fastest evolution in PB 8640M as was observed also in

injection-molded specimens (Fig. 16). Nevertheless, the peculiar increase at annealing temperature of -22 °C

is observed in all the materials; in respect to overall faster evolution in PB 8640M this increase occurs earlier

than in homopolymers. This fact evidences that it is not an accidental error as one could suggested.

Similarly, at annealing temperature of 60 °C the transformation rate is the slowest with distinct plato.

When evolutions in both injection-molded and extruded specimens (Figs 14-16 and 24-26) are

compared distinct similarities and differences resulting from different processing technology can be seen.

(Note: different tensile testing apparatus were used and the tensile modulus in injection-molded specimens

was calculated using software function without extensometer, while in extrudates an extensometer was used

for calculation.) The similarities were found in evolution at annealing temperatures +5, +22 and +40 °C in all

materials, nevertheless values of tensile modulus - are remarkably higher for extrudates. Probably differences

in injection-molded morphology (skin-core), where typical oriented structure in the skin can be observed

(created upon faster cooling), while in extrudates un-oriented structure was observed (Fig. 23), can play an

important role in transformation rate and affects final mechanical properties. In all materials at temperature

of +60 °C the evolution seems to possess similar trend in the evolution.

On the other hand, the modulus evolution at temperature of -22 is completely different except PB

8640M. A peculiar increase in injection-molded specimens at this temperature can be partially observed

remarkably later in time. This observation can confirm significant influence of processing and resulting

morphology on phase transformation.

It worth noting that both sets of specimens (injection molded and extruded) were also visually

analysed immediately after processing - actually, it was possible to stretch the injection-molded specimens

readily by hand. However, even although both sets of specimens possess approx. the same cross-section

55

areas, it was not possible to stretch the extruded specimens. This confirms the remarkable differences in

tensile modulus when comparing both specimens.

Fig. 24 Evolution of tensile modulus in extruded PB 0300M

Fig. 25 Evolution of tensile modulus in extruded DP 0401M

0 10 20 30 40 50 60 70 80 90

150

200

250

300

350

400

450

500

550

600

tens

ile m

odul

us (M

Pa)

annealing time (days)

-22 °C +5 °C +22 °C +40 °C +60 °C

0 10 20 30 40 50 60 70 80 90

150

200

250

300

350

400

450

500

tens

ile m

odul

us (M

Pa)

annealing time (days)

-22 °C +22 °C +60 °C

56

Fig. 26 Evolution of tensile modulus in extruded PB 8640M

5.3.4 Tensile elongation at break

Again, the evolution of tensile elongation at break illustrated in Figs 27-29 shows decreasing trend

compared to increasing tensile modulus. What can be seen is remarkably different evolution for each

material. In PB 8640M (Fig. 29) no distinct trend can be observed, it can be probably caused by extremely

high elongation – 400 ~ 540 %. On the other hand, in homopolymers trends in evolution are obvious and

even they differ. In both homopolymers (Figs 27 and 28) at annealing temperature of +60 °C the evolution

possesses virtually the same trend as well as at +22 °C. When annealing temperature of -22 °C is compared

in both Figs, in case of PB 0300M significant drop is observed between day 14 and 22, and further annealing

shows plateau in evolution. While in DP 0401M the elongation at break is gradually decreasing but not fully

correlating with tensile modulus (Fig. 25).

When extruded and injection-molded (Figs 18 and 19) results of elongation at break are compared,

remarkable differences are clearly seen. In case of extrudates PB 0300M the elongation is 3 ~ 4 times higher

than as injection-molded specimens, on the other hand in copolymer PB 8640M the elongation is even 9 ~ 10

times higher; in case of injection-molded DP 0401M the elongations were over 120 % and out of testing

possibility on employed testing apparatus. These elongation results certainly confirms the effect of

processing on the phase transformation respectively on final mechanical properties.

0 10 20 30 40 50 60 70 80 9080

100

120

140

160

180

200

220

240

260

280

300

320

340

360

tens

ile m

odul

us (M

Pa)

annealing time (days)

-22 °C +5 °C +22 °C +40 °C +60 °C

57

Fig. 27 Evolution of tensile elongation at break in extruded PB 0300M

Fig. 28 Evolution of tensile elongation at break in extruded DP 0401M

0 10 20 30 40 50 60 70 80 90220

240

260

280

300

320

340

360

380

400

420

440

tens

ile e

long

atio

n at

bre

ak (%

)

annealing time (days)

-22 °C +5 °C +22 °C +40 °C +60 °C

0 10 20 30 40 50 60 70 80 9080

100120140160180200220240260280300320340360380400420

tens

ile e

long

atio

n at

bre

ak (%

)

annealing time (days)

-22 °C +22 °C +60 °C

58

Fig. 29 Evolution of tensile elongation at break in extruded PB 8640M

5.3.5 Tensile strength at break

Following Figs 30-32 illustrate evolution of tensile strength at break, which particularly corresponds

to evolution of tensile elongation at break Figs 27-29. Again, no reliable differences of evolution in

PB 8640M were observed; on the contrary, in homopolymers significant decrease of strength at break can be

seen.

However, these results possess an opposite trend upon annealing time with those observed in

injection-molded specimens (Figs 20 and 21). The strengths at break are in case of PB 0300M quite

equivalent, except temperatures of -22 and +60 °C; but in PB 8640M the strengths are significantly higher

for the extrudates. Generally, one would expect higher drawability and possibly strength at break in

injection-molded specimens possessing lower modulus. On the other hand, an enhanced molecular

orientation in injection-molded specimens can reduce in general the drawability. However, this rather

surprising finding would still require further to be elucidated.

0 10 20 30 40 50 60 70 80 90380

400

420

440

460

480

500

520

540

tens

ile e

long

atio

n at

bre

ak (%

)

annealing time (days)

-22 °C +5 °C +22 °C +40 °C +60 °C

59

Fig. 30 Evolution of tensile strength at break in extruded PB 0300M

Fig. 31 Evolution of tensile strength at break in extruded DP 0401M

0 10 20 30 40 50 60 70 80 9019202122232425262728293031323334

tens

ile s

treng

th a

t bre

ak (M

Pa)

annealing time (days)

-22 °C +5 °C +22 °C +40 °C +60 °C

0 10 20 30 40 50 60 70 80 9019202122232425262728293031323334

tens

ile s

treng

th a

t bre

ak (M

Pa)

annealing time (days)

-22 °C +22 °C +60 °C

60

Fig. 32 Evolution of tensile strength at break in extruded PB 8640M

0 10 20 30 40 50 60 70 80 9023

24

25

26

27

28

29

30

31

32

33

34

35

36

tens

ile s

treng

th a

t bre

ak (M

Pa)

annealing time (days)

-22 °C +5 °C +22 °C +40 °C +60 °C

61

6 MULTISCALE CHARACTERIZATION OF DEGRADATION IN PB-1 The degradation behavior of poly(1-butene) is not still clear yet; it supposed to be similar to isotactic

polypropylene due to presence of tertiary carbon on the backbone. The formation of peroxidic species in the

polymer backbone is encountered as the primary step of the oxidative degradation. The hydroperoxide

groups not only play an important role in inducing both thermal and photo-oxidation of polymers but they

are also important intermediates in the overall oxidation reactions. Thus, (hydro)peroxides are the key

products in the clear understanding of a mechanism, as well as to gain better insight into correlation between

chemical evolution and evolution of the molecular structure (scission/recombination/crosslinking) [5].

Past works related to the phase transformation behavior claims that it is influenced: by physical

ageing conditions as temperature, pressure and applied stress; by some additives; by molecular

characteristics, addition of ethylene copolymers. However, no papers deal with the influence of degradation

on the phase transformation behavior which should subsequently play a role in the possible

utilization/recycling of the material at the end of its lifetime.

To follow this idea, a multiscale characterization has been performed on two sorts of specimens. The

compression-molded specimens with thickness of 0.3 mm were used for study of degradation during early

stages 0 ~ 40 UV irradiation hours or 0 ~ 69 days of artificial weathering. The prolonged times led to

disintegration of such specimens. Therefore, the specimens with thickness of 1 mm were used for UV

irradiation times 0 ~ 160 hours, respectively 0 ~ 103 days naturally weathered.

6.1 Early stages

6.1.1 DSC study

The melting behavior of degraded specimens of both homopolymers (DP 0401M and PB 0300M) is

illustrated in Fig. 33. It can be seen that the melting temperature of the form I, observed by first melting

(which corresponds to melting of the form I), vary interval of 122 ~ 127 °C in both materials. No distinct

trend is observed compared to the second melting (performed immediately after first melting and controlled

crystallization), which corresponds to the melting of the form II. In both materials a decrease of melting

temperature with increasing irradiation time is observed with total drop of 10 ~ 12 °C between time 0 and 40

hours. Upon the transformation from form II to I, virtually identical thermal behavior with shifted melting

temperature can be seen (melting after transformation is completed corresponds to the melting of the form I

after controlled crystallization and annealing during which the transformation II – I occurs), as the slope of

m2 and mtr curve in Fig. 33 is similar.

What can be expected in the initial stage of degradation is a prevailing effect of chain scission, which

can positively influence the crystallization kinetics. As a consequence, final structure is thermodynamically

more stable, and its melting temperature in short-term irradiated specimens can even exceed that of non-

irradiated. This is particularly evident in PB 0300M from the melting of phase I upon transformation.

62

However, chain scission is followed by the introduction of impurities into the molecular structure. This effect

leads to the decrease of crystallizability and the crystallization to less stable structure [106]. In thermograms,

it is presented by gradual decrease of melting temperatures of both phase II and I with prolonged irradiation

time.

Fig. 33 Melting behavior of PB 0300M and DP 0401M (m1 – first melting, m2 – second melting, mtr – melting after transformation is completed)

Fig. 34 Crystallization behavior of PB 0300M and DP 0401M

0 5 10 15 20 25 30 35 40102

104

106

108

110

112

114

116

118

120

122

124

126

128

130

mel

ting

tem

pera

ture

(°C

)

UV irradiation time (hours)

m1 DP 0401M m1 PB 0300M m2 DP 0401M m2 PB 0300M mtr DP 0401M mtr PB 0300M

0 5 10 15 20 25 30 35 40

55

60

65

70

75

80

85

90

95

crys

talli

zatio

n te

mpe

ratu

re (

°C)

UV irrad iation time (hours)

DP0401M PB0300M

63

The effect of photo-induced changes on crystallization is even more evident from the crystallization

temperature (Tc), as shown in Fig. 34. In the case of PB 0300M, the non-monotonic evolution of Tc can be

seen. At the beginning, a significant increase of Tc is observed from 78 to 91 °C and after 10 hours of UV

irradiation the initial Tc is reached. This observation confirms that the chains scission positively influences

the crystallization rate (molecular mobility increases) and this effect prevails in the initial stage of UV

irradiation. However, an increasing number of introduced heterogeneities in polymer chains then causes

a linear drop of Tc about 30 °C between 5 and 20 UV irradiation hours. Then in time, a gradual decrease of Tc

to 55 °C occurs. The evolution in Tc of DP 0401M upon UV irradiation is slightly different. Crystallization

temperature decreases with exposure time. During first 5 hours, degradation leads into a significant drop of

Tc, then, from 15 up to 40 hours only a small decrease can be seen up to 55 °C, i.e. the same values as that of

PB 0300M. Clearly, the first stage with increasing crystallization rate upon chain scission is missing. It can

be supposed that the competition between (i) the increase in crystallization kinetics upon chain scission and

(ii) the retardation of crystallization caused by introduction of molecular irregularities is influenced by initial

molecular weight. Indeed, the feasible maximum of crystallization rate is faster reached in lower-molecular

systems, thus in DP 0401M, in this case its crystallization temperature decreases even upon the initial stage

of UV irradiation.

In order to compare a thermal behavior of both specimens artificially and naturally degraded, the

material DP 0401 was exposed to natural weathering. Subsequent melting and crystallization temperatures

are illustrated in Fig. 35.

Fig. 35 Melting and crystallization behavior of DP 0401M exposed to natural weathering, (m1 – first melting, m2 – second melting, mtr – melting after transformation is completed)

0 5 10 15 20 25 30 35 40 45 50 55 60 65

60

70

80

90

105

110

115

120

125

130 m1 m2 mtr crystallization

exposure time (days)

tem

pera

ture

(°C

)

64

Similarly to accelerated weathering tests, the melting temperature of form I upon natural test does

not significantly changed. Upon the recrystallization, melting temperatures of both form I and II firstly

increases and then drops to low values, again similarly to changes arising from accelerated weathering.

However, the crystallization temperature shows non-monotonic evolution, which cannot be found within

accelerated ageing. Because of weaker intensity and different spectra of UV light are applied during natural

weathering, the molecular degradation is slowed down which is reflected by rather similar crystallization

behavior as in the case of PB 0300M upon artificial weathering. Upon 18 days of exposition, an increase of

Tc is manifested, then, between the days 18 and 28 a significant drop in Tc can be seen. Further, no

significant changes occurred from day 50 to 63 in crystallization behavior. It should be noted that during

natural weathering and accelerated UV irradiation the final melting and crystallization temperatures are

similar.

The evolution of subsequent phase transformation in re-melted UV irradiated specimens is illustrated

in following Figs 36-41. The Fig. 36 shows evolution of form I content during phase transformation of

degraded DP 0401M. Fig. 37 illustrates the ratios of the melting peaks the forms I and II. From both Figs it is

evident that after two hours of UV irradiation no significant change in transformation rate is observed.

However, the change of transformation rate is evident between 3 and 10 irradiation hours, when decreasing

trend occurs which correlates to decrease of crystallization temperature (Fig. 34). Further minor change in

transformation rate between 15 and 40 irradiation hours confirms its dependence on crystallization

temperature

Fig. 36 Evolution of form I content during transformation of degraded specimens for various irradiation times in DP 0401M

0 1 2 3 4 5 6

0

10

20

30

40

50

60

70

80

90

100

cont

ent o

f for

m I

(%)

transformation time (days)

0 hour 1 hour 2 hours 3 hours 5 hours 10 hours 15 hours 20 hours 30 hours 40 hours

65

Fig. 37 Evolution of ratio of form I to II during transformation of degraded specimens for various irradiation times in DP 0401M

This fact is supported by Figs 38 and 39 in case of material PB 0300M. The transformation rate of

non-irradiated specimen which has a lower Tc compared to UV irradiation time of 1.5, 3 and 5 hours

possesses also slower transformation rate. Further irradiation (10 ~ 40 hours) is followed by decreasing trend

the same as in Fig. 34. Again, the competition between the molecular mobility and regularity is a key factor

influencing the transformation.

Fig. 38 Evolution of form I content during transformation of degraded specimens for various irradiation times in PB 0300M

1 2 3 40.1

1

10

100

10 hours 15 hours 20 hours 30 hours 40 hours

ratio

of f

orm

s I:I

I

transformation time (days)

0 hour 1 hour 2 hours 3 hours 5 hours

0 1 2 3 4 5

0

10

20

30

40

50

60

70

80

90

100

cont

ent o

f for

m I

(%)

transformation time (days)

0 hour 1.5 hours 3 hours 5 hours 10 hours 15 hours 20 hours 30 hours 40 hours

66

Fig. 39 Evolution of ratio of form I to II during transformation of degraded specimens for various irradiation times in PB 0300M

A virtually the same behavior can be seen in Figs 40 and 41, where the transformation rates also

follows the changes of crystallization temperatures in naturally exposed specimens.

Fig. 40 Evolution of form I content during transformation of degraded specimens for various exposition times in DP 0401M

1 2 3 4 50.1

1

10

100

10 hours 15 hours 20 hours 30 hours 40 hours

ratio

of f

orm

s I:I

I

transformation time (days)

0 hour 1.5 hours 3 hours 5 hours

0 1 2 3 4 5 6 7

0

10

20

30

40

50

60

70

80

90

100

cont

ent o

f for

m I

(%)

transformation time (days)

0 day 18 days 29 days 50 days 57 days 63 days

67

Fig. 41 Evolution of ratio of form I to II during transformation of degraded specimens for various exposition times in DP 0401M specimens

When comparing Figs 37, 39 and 41 it can be seen that curves of the ratios of the forms I and II are

particularly superposed to each other. Thus, the transformation rate follows the changes in crystallization

temperature; the process is overall slower but with similar mechanism.

6.1.2 FTIR characterization

UV irradiation leads to significant changes in the IR absorption spectra in PB-1. The evolution of

chemical species in degraded PB-1 can be seen in following Figs 42-44. The major by-products resulting

from presence of atmosphere oxygen are carbonyl products. As the oxidation proceeds, broad absorption

bands growth at: 1640, 1712, 1735 and 1780 cm-1 are observed. A gradual formation of carboxylic acid

(1712 cm-1) can be seen in all cases. With further exposition the formation of esters becomes dominant at

1735 cm-1 in case of DP 0401M; the formation at absorption of 1780 cm-1 is assigned to peresters or

γ-lactones; a weak peak at absorption 1640 cm-1 is assigned to vinyl groups [107].

The UV irradiation in DP 0401M and PB 0300M (Figs 42 and 43) shows similar mechanism of

degradation, however in DP 0401M the evolution of chemical species is quicker in the initial stage compared

to PB 0300M, and as written above the esters (1735 cm-1) become dominant from 60 hours. This behavior

follows the DSC observation; the interrelation with the physical effect of transformation is obvious.

1.0 1.5 2.0 2.5 3.0

0.1

1

10

100

ratio

of f

orm

s I:I

I

transformation time (days)

0 day 18 days 29 days 50 days 57 days 63 days

68

Fig. 42 FTIR spectra of DP 0401M - evolution in carbonyl area during UV irradition

Fig. 43 FTIR spectra of PB 0300M - evolution in carbonyl area during UV irradiation

0,1

0,2

0,3

0,4

0,5

0,6

0,7

0,8

0,9

1,0

1,1

1,2

Abs

orba

nce

1650 1700 1750 1800 1850 Wavenumbers (cm-1)

0,2

0,3

0,4

0,5

0,6

0,7

0,8

0,9

Abs

orba

nce

1650 1700 1750 1800 1850 Wavenumbers (cm-1)

UV

irra

diat

ion

time

0, 1

0, 2

0, 3

0, 4

0, 5

0, 6

0, 7

0 ho

urs

UV

irra

diat

ion

time

0, 1

0, 2

0, 3

0, 4

0, 5

0, 6

0, 7

0 ho

urs

69

Fig. 44 FTIR spectra of DP 0401M - evolution in carbonyl area during natural exposition

On the other hand comparing UV irradiation and natural exposition in DP 0401M (Figs 42 and 44)

shows significantly different behavior where carbonyl groups (1712 cm-1) are predominant.

When absorbancies at 1712 cm-1 are plotted versus UV irradiation time or exposition time, as can be

seen in Fig. 45, the chemical evolution of PB 0300M is weaker than in DP 0401M. Comparing both

degraded DP 0401M it can be seen that natural weathering lead to similar chemical evolution during first 29

days what is in accordance with 30 hours of UV irradiation. Nevertheless, further natural exposition supports

formation of carboxylic acids groups (Fig. 42) while in UV irradiated specimens esters groups compete with

carboxylic acids.

It should also be noted that FTIR is not able to detect peroxide species, which are during radical

reactions transformed to detectable chemical groups. Nevertheless, several reliable methods exist, such as

modulated DSC observation, iodometric titration or indirect method based e.g. on rheological behavior.

0,2

0,3

0,4

0,5

0,6

0,7

0,8

0,9

1,0

1,1

1,2

1,3

Abs

orba

nce

1650 1700 1750 1800 1850

Wavenumbers (cm-1)

expo

sitio

n tim

e

0, 1

8, 2

9, 4

3, 4

9, 6

5 da

ys

70

Fig. 45 Comparison of absorbancies at 1712 cm-1 in UV irradiated and naturally exposed specimens

6.1.3 Rheology

Figs 46 and 47 show the Cole-Cole plots for the individual times of UV irradiation. In the complex

plane this model predicts the variation of the viscosity components (η” versus η’) to be an arc of circle [104].

From this plot it is easy to determine zero shear viscosity η0 which depends on the molecular weight and

follows a power law. It is evident that this is convenient method to determine molecular changes in the short

UV irradiation times for PB-1 with respect to MW. In case of DP 0401M (Fig. 46) at the beginning of UV

irradiation (1.6 hour) η0 slightly decreases. During further UV irradiation, 5 hours, η0 drops and then further

significant drop is observed after 10 hours, which shows practically immeasurable low viscosity on

employed apparatus, respectively chain scissions occur. Material PB 0300M (Fig. 47) shows at the beginning

of UV irradiation virtually the same behavior as DP 0401M. On the other hand during further UV irradiation

the decrease of η0 is progressive which is contributed to approx. 3 times higher η0 (higher MW) than DP

0401M. It is obvious that chain scission reactions, initiated by hydroperoxides decomposition during melt

rheological measurements, are dominant during initial stage of UV irradiation even if recombination

reactions occur [75].

0 5 10 15 20 25 30 35 40 45 50 55 60 65 70

0.0

0.1

0.2

0.3

0.4

0.5

0.6

0.7

0.8

0.9

1.0

1.1

0 10 20 30 40 50 60 70

abso

rban

ce a

t 171

2 cm

-1

UV irradiation time (hours)

UV irradiated DP 0401M UV irradiated PB 0300M

naturally exposed DP 0401M

exposition time (days)

71

Fig. 46 Cole-Cole plot of DP 0401M up to 10 hours of UV irradiation

Fig. 47 Cole-Cole plot of PB 0300M up to 10 hours of UV irradiation

0 500 1000 1500 2000 2500 30000

100

200

300

400

500

600

700

800

900

1000

1100

1200

η" (P

a.s)

η' (Pa.s)

0 hours UV 1.6 hours UV 3.3 hours UV 5 hours UV 10 hours UV

0 2000 4000 6000 8000 100000

500

1000

1500

2000

2500

3000

3500

4000

4500

η" (P

a.s)

η' (Pa.s)

0 hours UV 1.6 hours UV 4 hours UV 6.6 hours UV 10 hours UV

72

The dependences of extrapolated zero shear viscosity η0 and distribution parameter h on UV

irradiation time are illustrated in Fig. 48. The decrease of η0 directly reflects changes in molecular weight,

while parameter h reflects the changes in molecular weight distribution. The drop of h at time of 10 hours

signifies double distribution, which can be particularly seen in Fig. 47.

It is evident that DP 0401M degrades more readily because of lower molecular weight compared to

PB 0300M which is consistent with observation on the phase transformation.

Fig. 48 Evolution of zero shear viscosity and distribution parameter during UV irradiation

0 1 2 3 4 5 6 7 8 9 100

500100015002000250030003500

7000

8000

9000

10000

11000

UV irradiation time (hours)

η 0 (P

a.s)

DP0401M PB0300M

0 1 2 3 4 5 6 7 8 9 10

0.38

0.39

0.40

0.41

0.42

0.43

0.44

0.45

0.46

0.47

0.48

dist

ribut

ion

para

met

er h

UV irradiation time

DP0401M PB0300M

73

6.2 Prolonged stages

6.2.1 DSC study

The melting behavior of UV irradiated 1 mm thick sheets of DP 0401M is illustrated in Fig. 49. The

similar trend of decreasing melting and crystallization temperatures were observed in Figs 33 and 34 for the

sheet with thickness of 0.3 mm which allowed detail insight to initial stage of degradation in PB-1.

Fig. 49 Melting and crystallization behavior of PB 0401M (m1 – first melting, m2 – second melting, mtr – melting after transformation is completed)

6.2.2 FTIR-ATR characterization

Similarly, the comparing of FTIR and FTIR-ATR spectra of DP 0401M (Figs 42 and 50) shows

virtually the same trend in the beginning of UV irradiation. No changes in first 10 hours, then time of 20

hours an increase in carboxylic area (1712 cm-1). Since 60 hours of UV irradiation, a prevailing growth of

esters occurs (1735 cm-1).

0 20 40 60 80 100 120 140 160

55

60

65

70

105

110

115

120

125

m1 m2 mtr crystallization

irradiation time (hours)

tem

pera

ture

(°C

)

74

Fig. 50 FTIR-ATR spectra of DP 0401M - evolution of carbonyl area during UV irradiation

6.2.3 Rheology

The thicker sheets also enables to perform rheology experiments in wider time range, because of the

penetration of hydroperoxides groups is more gradual in thickness profile of 1.0 mm compared to 0.3 mm.

The rheological measurements of naturally exposed specimens of DP 0401M were practically

valueless, because after of only 18 days of exposition, the value of zero shear viscosity η0 dropped to

<100 Pa.s.

The rheological behaviors represented by Cole-Cole plots of both materials are illustrated in Figs 51

and 52. In both materials, UV irradiation lead to virtually the same decrease of zero shear viscosity η0 after

15 ~ 20 hours. Nevertheless, the development of degradation is completely different during initial 10 hours.

In DP 0401M (Figs 51 and 53) during first 2.5 hours a recombination reactions take place in

competition with chain scission reactions or the effect of chemi-crystallization. Thereafter, gradual decrease

of η0 is observed, caused by chain scission reactions and increasing amount of hydroperoxides groups. The

changes in parameter h were not observed.

0,02

0,04

0,06

0,08

0,10

0,12

0,14

0,16

0,18

0,20

0,22

Abs

orba

nce

1600 1700 1800 1900 Wavenumbers (cm-1)

UV

irra

diat

ion

time

0, 1

0, 2

0, 3

0, 4

0, 6

0, 1

20, 1

60

75

Fig. 51 Cole-Cole plot of DP 0401M up to 20 hours of UV irradiation

Fig. 52 Cole-Cole plot of PB 0300M up to 20 hours of UV irradiation

The Cole-Cole plot of PB 0300M (Fig. 52) shows a particularly the same behavior as in the Fig. 47

for initial stage. Even if no significant changes in η0 were observed, the parameter shows its increase as in

the Fig. 48. After 10 hours of UV irradiation η0 remarkably drops to half-value of that at 7.5 hours.

Continuous UV irradiation led to further drops in η0 similarly as in DP 0401M. The changes of parameter h

are illustrated in Fig. 54, a similar increasing trend was observed as for sheets with thickness of 0.3 mm (see

Fig. 48).

0 500 1000 1500 2000 2500 3000 3500 4000 4500

0

200

400

600

800

1000

1200

1400

1600

1800

η" (P

a.s)

η' (Pa.s)

0 hours UV 2.5 hours UV 5 hours UV 7.5 hours UV 10 hours UV 15 hours UV 20 hours UV

0 2000 4000 6000 8000 10000 12000 14000

0

500

1000

1500

2000

2500

3000

3500

4000

4500

5000

5500

6000

η" (P

a.s)

η' (Pa.s)

0 hours UV 5 hours UV 7.5 hours UV 10 hours UV 15 hours UV 20 hours UV

76

Fig. 53 Evolution of zero shear viscosity and distribution parameter during UV irradiation in DP 0401M

Fig. 54 Evolution of zero shear viscosity and distribution parameter during UV irradiation in PB 0300M

0 2 4 6 8 10 12 14 16 18 200

500

1000

1500

2000

2500

3000

3500

4000

4500

η0

h

UV irradiation time (hours)

zero

she

ar v

isco

sity

η0 (P

a.s)

0.390

0.395

0.400

0.405

0.410

0.415

0.420

0.425

0.430

0.435

0.440

dist

ribut

ion

para

met

er h

0 2 4 6 8 10 12 14 16 18 200

200400600800

1000120014006000

7000

8000

9000

10000

11000

12000

13000

14000

15000

UV irradiation time (hours)

zero

she

ar v

isco

sity

η0 (

Pa.s

)

0.370

0.375

0.380

0.385

0.390

0.395

0.400

0.405

0.410

0.415

0.420

η0

h

dist

ribut

ion

pera

met

er h

77

6.2.4 WAXS characterization

In the Figs 55 and 56 X-ray data, namely crystallinity and form II content of both sets of degraded

specimens of material DP 0401M can be seen. It is obvious that no distinct trends were observed and both

calculated values vary within close interval. This is actually with agreement with previously presented DSC

data where the first melting varies within close interval for both types of degradation.

Fig. 55 Evolution of crystallinity and form II content in UV irradiated specimens DP 0401M

Fig. 56 Evolution of crystallinity and form II content in naturally exposed specimens DP 0401M

0 10 20 30 40 50 60 70 80 90 100 110 120 130 140 150 16050

55

60

65

70

75

80

0

2

4

6

8

10

12

14

16

18

20

crys

talli

nity

(%)

UV irradiation time (hours)

crystallinity form II content

form

II c

onte

nt (%

)

0 10 20 30 40 50 60 70 80 90 100 110 120 130 14050

55

60

65

70

75

80

crystallinity form II content

exposition time (days)

crys

talli

nity

(%)

0

2

4

6

8

10

12

14

16

18

20

form

II c

onte

nt (%

)

6.2.5 Ste

In o

was used to

surface crac

deep crack

create a netw

network by

6.2.6 Sca

A d

59). As can

from center

differences

cracks; this

individual l

reomicrosco

order to obse

o take image

cks. Since 40

lines after 6

work. Furthe

widening th

Fig. 57 I

anning Elect

deeper look i

n be seen in

r of spherulit

are visible.

s could cau

amellae.

opy

erve the macr

s of the degr

0 hours of UV

0 hours. As

er degradatio

he crack lines

Images of sur

tron Micros

inside the str

Fig. 58 the

tes. For comp

The similar

sed by diffe

roscopic effe

raded specim

V irradiation

for 80 hours

on in time of

s and penetra

face degradat

copy

ructure was

cracks arise

parison, all t

morphology

ferent MW a

ect of UV irr

mens (see Fig

n a tiny crack

s, it can be s

f 120 and 160

ating into dep

ion in DP 040

performed u

e on spheruli

three materia

y of PB 0300

and probabl

radiation on

g. 57). What

k nucleus are

seen that cra

0 hours lead

pth.

01M in variou

using scannin

ite level follo

als were UV

0M does not

ly higher nu

surface chan

t is obvious

observed, co

ack lines inte

to propagati

us UV irradiati

ng electron m

owing the la

irradiated fo

t enable such

umber of tie

nges, a stereo

is the forma

onsequently

erconnect eac

ion of such c

ion times

microscopy (

amellar grow

or 160 hours

h intensive f

e molecules

78

omicroscope

ation of deep

enlarging to

ch other and

created crack

(Figs 58 and

wth direction

and distinct

formation of

s connecting

8

e

p

o

d

k

d

n

t

f

g

79

Fig. 58 SEM images of surface degradation in the materials for various UV irradiation times

80

While Fig. 58 shows surfaces of UV accelerated irradiated specimens, in Fig. 59 the surfaces of

naturally degraded specimens can be seen. The formation of cracks is different and especially high amount of

smaller cracks without distinct evolution, which is preferentially caused by weathering and erosions.

Fig. 59 SEM images of surface degradation in the DP 0401 naturally degraded at various exposition times

81

6.3 Degradation of extruded PB-1 In previous Chapters 6.1 and 6.2 the compression molded sheets were analyzed using various

experimental techniques regarding their degradation and subsequent phase evolution. On the other hand, the

following subchapter studies the effect of photodegradation on extruded specimens allowing also the

assessment of mechanical properties. Two degradation regimes were used; first set of specimens was UV

irradiated from 0 to 300 hours, seconds set of specimens was exposed to natural weathering in period from

June 6 to September 17 2008 e.g. 0 ~ 103 days. Then, the degraded sets of specimens were characterized by

WAXS, FTIR-ATR, stereo-microscopy and mechanical testing.

6.3.1 Mechanical properties

The evolutions of mechanical properties in DP 0401M during UV irradiation and comparison with

natural weathering are illustrated in following Figs 60-62. It is obvious that the natural weathering led to

different results in all three measured parameters compared to accelerated UV irradiation. The evolution of

tensile modulus (see Fig. 60) is followed in the case of accelerated UV irradiation by a gradual decreased

from approx. 500 to 350 MPa, except the variation in the initial time 0 ~ 20 hours. On the contrary, natural

weathering lead to a slightly increasing trend of modulus; actually these values partly correspond to the UV

irradiation time of 20 hours. This can be related to the chemi-crystallization phenomenon.

Fig. 60 Evolution of tensile modulus in UV irradiated and naturally exposed specimens

The Fig. 61 illustrates the tensile strength at break behavior of degraded specimens. Similarly as in

the tensile modulus the UV irradiation time between 0 ~ 20 hours shows a specific behavior with a slight

increase (or no change) in 10 hours being probably caused by mentioned chemi-crystallization, but the

further irradiation lead to drop from 21 to 15 MPa . From 40 hours, the further UV irradiation causes gradual

decrease to 11.5 MPa. On the other hand, the behavior of natural weathered specimens leads to

0 10 20 30 40 50 60 70 80 90

340

360

380

400

420

440

460

480

500

520

540

560

580

600

0 20 40 60 80 100 120 140 160 180 200 220 240 260 280 300

340

360

380

400

420

440

460

480

500

520

540

560

580

600

exposition time (days)

tens

ile m

odul

us (M

Pa)

UV irradiation time (hours)

82

an insignificant decrease of strength at break after 24 days of exposition. Nevertheless, after 35 days a drop

to a half of initial value occurred, followed by a decrease to the values between 8 ~ 9 MPa. Thus the

determination of degradation factor between both, accelerated UV irradiation and natural weathering is

practically restricted to the short exposition times to 35 days, and 40 hours respectively.

Fig. 61 Evolution of tensile stress at break in UV irradiated and naturally exposed specimens

Again, the specific evolution in the initial UV irradiation time of 0 ~ 40 hours can be seen in Fig. 62.

The elongation at break significantly drops from approx. 38 % to a few percent after 20 hours. Further

irradiation does not influences elongation at break and the value vary between 9 ~ 11 %; it is worth noting

that such values are already rather low, probably on their minima. Similar behavior is observed in naturally

exposed specimens, after the initial drop to 10 %, the elongation vary in interval 3 ~ 5 %. It can be suggested

that 20 ~ 40 hours of UV irradiation corresponds to 24 days of natural weathering, nevertheless this factor

cannot be applied for longer exposition time. It should be mentioned that the both effect the chemi-

crystallization and chain scission will in general lead to a decrease of the material drawability.

Fig. 62 Evolution of tensile elongation at break in UV irradiated and naturally exposed specimens

0 20 40 60 80 100 120 140 160 180 200 220 240 260 280 300789

10111213141516171819202122

0 10 20 30 40 50 60 70 80 90789

10111213141516171819202122

tens

ile s

treng

th a

t bre

ak (M

Pa)

UV irradiation time (hours)

exposition time (days)

0 20 40 60 80 100 120 140 160 180 200 220 240 260 280 3000

5

10

15

20

25

30

35

40

45

0 10 20 30 40 50 60 70 80 900

5

10

15

20

25

30

35

40

45

tens

ile e

long

atio

n at

bre

ak (%

)

UV irradiation time (hours) exposition time (days)

83

The testing of mechanical properties offers a suggestion that natural exposition with lower UV

intensity leads to the chemi-crystallization, which slightly increases tensile modulus. Therefore, naturally

exposed specimens become more brittle as has been observed in its tensile elongation at break and strength at

break. On the contrary, in the UV irradiated specimens the chemi-crystallization process and modulus

increase is overshadowed by the erosion of crystallites themselves upon strong accelerated UV irraditation.

6.3.2 FTIR-ATR characterization

FTIR-ATR characterization of degraded specimens was performed to understand the evolution of

mechanical properties. From comparison of Figs 63 and 64 it is obvious that different chemical evolution

occurred. In UV irradiated specimens (Fig. 63), a unvalued difference between time 0 and 10 hours can be

seen within region 1700 ~ 1750 cm-1, and with rising time 20 and 40 hours gradual increase is observed.

Further irradiation lead to gradual increase in this region. The increase of peak in region 1645 cm-1 which

belongs to vinyl groups is observed in both sets of degraded specimens. On the contrary, in the naturally

exposed specimens (Fig. 64) within region 1700 ~ 1750 cm-1 the significant step is observed after 24 days

and the formation of carboxylic acids (1712 cm-1) is pronounced as well in all naturally exposed specimens.

When absorbance values are compared (see Fig. 65), UV irradiation time of 20 hours corresponds to

the exposition times 24 ~ 37 days and further exposition days lies between 20 ~ 40 hours.

From Fig. 65 the chemi-crystallization influencing mechanical properties can be clarified. In

naturally exposed specimens positive effect of chemi-crystallization has lead to increase of modulus which is

correlates to moderate increase of carboxylic acids with partial plato in the surface layer. On the other hand,

in UV irradiated specimens the content of chemical groups sharply increase resulting probably also in

crystallite erosion and material disintegration.

84

Fig. 63 FTIR-ATR spectra of UV irradiated specimens

Fig. 64 FTIR-ATR spectra of naturally exposed specimens

0,02

0,03

0,04

0,05

0,06

0,07

0,08

0,09

Abs

orba

nce

1550 1600 1650 1700 1750 1800 1850 Wavenumbers (cm-1)

0,020

0,025

0,030

0,035

0,040

0,045

0,050

Abs

orba

nce

1600 1700 1800

Wavenumbers (cm-1)

expo

sitio

n tim

e

0, 2

4, 3

7, 5

8, 7

2, 8

8, 1

02 d

ays

UV

irra

diat

ion

time

0, 1

0, 2

0, 4

0, 6

0, 8

0, 1

00,

120,

160

, 200

, 240

, 300

hou

rs

85

Fig. 65 Comparison of absorbancies at 1712 cm-1 in UV irradiated and naturally exposed specimens

6.3.3 WAXS characterization

X-ray scattering data with crystallinity and form II content of both sets of degraded specimens are

illustrated in Figs 66 and 67. It should be noted that both sets were prepared from the same extruded tape

however the WAXS was performed in reflection mode and improvement of crystallinity in the core may not

be observed, although the most obvious effect of photodegradation can be anyhow expected in the skin-layer.

It can be seen in both Figs that crystallinity vary within interval of approx. 2 % as has been observed in

previous chapters and first melting in DSC study. What is important is the fact that crystallinity of naturally

exposed specimens is slightly higher than in UV irradiated specimens, and do not virtually decrease during

degradation. In both Figs, distinct changes in the evolution of form II content can be seen. Upon UV

irradiation form II content of UV irradiated specimens quite fluently decreases from approx. 6 to 3 %. On the

contrary, in naturally exposed specimens different variation of form II content can be seen. The slight

increase can be probably caused by chemi-crystallization, however, further exposition lead to the decrease to

the virtually the same from II content as in the UV irradiated specimens.

0.00

0.01

0.02

0.03

0.04

0.05

0.06

0.07

0.08

0 20 40 60 80 100 120 140 160 180 200 220 240 260 280 300

0 10 20 30 40 50 60 70 80 90

UV irradiated samples

abso

rban

ce a

t 171

2 cm

-1

exposition time (days)

naturally exposed samples

UV irradiation time (hours)

86

Fig. 66 Evolution of crystallinity and form II content in UV irradiated specimens

Fig. 67 Evolution of crystallinity and form II content in naturally exposed specimens

0 50 100 150 200 250 30050

55

60

65

70

75

80

crystallinity

UV irradiation time (hours)

crys

talli

nity

(%)

0

1

2

3

4

5

6

7

8

9

10

form II content

form

II c

onte

nt (%

)

0 10 20 30 40 50 60 70 80 90 10055

60

65

70

75

80

85

crystallinity form II content

exposition time (days)

crys

talli

nity

(%)

0

1

2

3

4

5

6

7

8

9

10

form

II c

onte

nt (%

)

6.3.4 Ste

The

Similarly a

compression

the crack lin

cracks are p

to widening

reomicrosco

e observation

s in the Fig

n molded sp

nes are in pa

perpendicular

g the crack lin

Fig.

opy

n of the deg

g. 57, initial

ecimens the

arallel direct

rly interconn

nes.

68 Images of

graded surfa

visible crac

initial crack

tion as extru

nects to a cra

f extruded tape

aces of the

cks appear a

ks possess ran

usion directio

ack network a

es DP 0401M

UV irradiate

after 40 hou

ndom directi

on. With incr

after 120 hou

M in various UV

ed specimen

urs of UV ir

ion of growth

reasing UV

urs; further U

V irradiation t

ns is shown

rradiation. W

h, in extrude

irradiation ti

UV irradiatio

times

87

in Fig. 68.

While in the

ed specimens

ime, parallel

on lead again

7

.

e

s

l

n

88

6.3.5 Effect of temperature on crystallinity and form II content

The effect of synergy of the temperature during UV irradiation was shortly studied using WAXS.

Thus, two series of specimens, which were stored at room conditions for 6 months, were used. First series of

them was UV irradiated at temperature of 60 °C and the other was annealed in oven with constant

temperature of 60 °C. Crystallinity and content of form II were calculated from WAXS data and plotted in

Figs 69 and 70.

It can be seen in Fig. 69 that crystallinity in both types of DP 0401M slightly increases in

dependence of time. This increase is more significant for annealing in oven at 60 °C, while during a synergy

of UV irradiation the crystallinity increase is lower indicating that the simultaneous processes of morphology

perfection are slow down, however just after long-time exposure. The difference however is rather small

(statistically insignificant) to openly talk about any real effects. Interestingly, further research of such

interrelations would be relevant (note in Fig. 55 the crystallinity of the UV degraded tapes of DP 0401M

varied and decreases with UV irradiation time).

Fig. 69 Evolution of crystallinity in annealed and degraded specimens

0 20 40 60 80 10065

70

75

80

85

crys

talli

nity

(%)

time (hours)

annealed sheet at 60 °C annealed tape at 60 °C degraded sheet at 60 °C degraded tape at 60 °C

89

Similarly, in Fig. 70 the evolution of the form II content shows virtually no significant differences

which could however changed upon long-term exposure. However, the decrease of crystallinity and form II

content reported in Fig. 66 can be then related mainly the effect of degradation.

Fig. 70 Evolution of form II content in annealed and degraded specimens

0 20 40 60 80 1000

1

2

3

4

5

6

7

8

9

10

form

II c

onte

nt (%

)

time (hours)

annealed sheet at 60 °C annealed tape at 60 °C degraded sheet at 60 °C degraded tape at 60 °C

90

7 CONCLUSION AND PERSPECTIVES In the bibliographic study, the present state of knowledge concerning the phase transformation and

degradation of PB-1 are introduced and discussed. The experimental part of this thesis has been then divided

into the two main parts: 1. a study of the phase transformation PB-1 and subsequent property evolution and

2. the role of UV light in such transformation.

The first part studies the influence of the temperature on the phase transformation from several sides

on various three PB-1 materials. Firstly, the DSC experiments were performed to observe thermal behavior

during annealing under five various temperatures. It was confirmed that addition of ethylene comonomer

positively increases molecular mobility and consequently the phase transformation. The fastest rate was

observed at annealing temperatures of +5 and +22 °C, while with increasing temperature up to 60 °C the

transformation slows down and the slowest rate occurred at -22 °C.

Then, the materials were processed by both injection-molding and extrusion technology to answer

the question whether processing with subsequent variation in morphology gradients can affect the phase

transformation being characterized by the evolution of tensile properties in the same materials and annealing

temperatures as was performed in DSC study. The DSC results were confirmed; however the evolution of

tensile properties on injection-molded or extrudates take remarkably longer time than in DSC specimens.

Probably thickness of the specimens can play a role in the rate of the phase transformation; further research

would be desirable. It was observed that the processing significantly influences elongation break and tensile

strength at break, which was higher in extrudates than in oriented injection-molded specimens.

The second part characterizing the degradation behavior of the material is actually the first study in

this area screening such effects for PB-1. Thermal analysis determined significant changes in crystallization

behavior influencing morphology arrangement and resulting thermal properties. Also, the effect of

degradation on the phase transformation has been proven; the degradation significantly slows down the

phase transformation. The significant differences were found between the accelerated UV irradiation and

natural weathering. Rheological behavior was suitable method for determination of early stages in

photodegradation of PB-1 using Cole-Cole plot, where other techniques do not properly record any relevant

changes. Microscopic methods showed the evolution on the surface of degraded specimens.

Further step was to study the effect of degradation on mechanical properties of the extrudates and

subsequent analysis to understanding. The effect of different design of the degradation procedure lead to

different evolution of tensile modulus. In naturally exposed specimens, the chemi-crystallization

phenomenon was observed resulting in increase of tensile modulus while in accelerated UV irradiated

specimens decrease of modulus was observed caused by erosion of crystallites in the surface layers. In both

sets, elongation at break and tensile strength decreases upon photodegradation. The photodegradation also

lead to a slight decrease of form II content in both sets.

91

Further promising research could be concerned on high-pressure behavior of PB-1, because as

written in bibliographic part, pressure increases the rate of the phase transformation however interrelations

with morphology has not been offered yet. Partial research on self-reinforcemement of PP and PB-1 blend

has already been done at the Department of Polymer Engineer. Another option based on a high PB-1

sensitivity to photodegradation could be an use of PB-1 as a suitable basic material for bio- (or oxo-)

degradable materials, even if the price is nowadays higher than for common polyolefins; on the other hand,

a proper stabilization could bring further applicability of PB-1 favorable based on its phase transformation.

92

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AUTHOR’S PUBLICATIONS Conference contributions

1. Beníček, L., Čermák, R., Verney, V., Commereuc, S., Obadal, M.; The effect of UV irradiation on melting and crystallization behaviour of poly(1-butene). Submitted for ANTEC 2009 - Society of Plastic Engineers, Chicago, IL, USA, June 22-24 2009

2. Hnidáková D., Čermák, R., Beníček, L. Navrátilová J., Vaysse, M.; The development of microfibrillar-phase morphology in PB-1/PP blends. Submitted for ANTEC 2009 - Society of Plastic Engineers, Chicago, IL, USA, June 22-24 2009

3. Beníček, L., Verney, V., Čermák, R., Commereuc, S.; Multiscale characterization of photodegradation in polybutene-1. Biannual Conference on Modification, Degradation and Stabilisation of Polymers, MoDeSt2008, Liege, Belgium, 2008

4. Beníček, L., Chvátalová, L., Čermák, R., Verney, V.; Evolution of mechanical properties during polymorphic transformation in polybutene-1. Biannual Conference on Modification, Degradation and Stabilisation of Polymers, MoDeSt2008, Liege, Belgium, 2008

5. Beníček, L., Verney, V., Čermák, R., Commereuc, S.; Characterization of photodegradation in polybutene-1: from early to prolonged stages. IUPAC 42th World Polymer Congress, MACRO 2008, Taipei, 2008

6. Beníček, L., Verney, V., Commereuc, S., Čermák, R., Mošnovská, R., Navrátilová, J.; Early stages of photodegradation in polybutene-1. Polymer Processing Society 24th Annual Meeting ~ PPS-24 ~ , Salerno, Italy, 2008

7. Beníček, L., Chvátalová, L., Čermák, R., Verney, V.; Polymorphic transformation in polybutene-1: temperature influence. Polymer Processing Society 24th Annual Meeting ~ PPS-24 ~ , Salerno, Italy, 2008

8. Beníček, L., Chvátalová, L., Čermák, R., Verney, V.; Thermal-induced transformation of structure and properties of polybutene-1. ANTEC 2008 - Society of Plastic Engineers, Milwaukee, WI, USA, 2008

9. Beníček, L., Verney, V., Commereuc, S., Čermák, R.; Characterization of photodegradation in polybutene-1 via rheological measurements. Journées des jeunes rhéologues 2008 (Groupe Français de Rhéologie), Erdeven, France, 2008

10. Beníček, L., Verney, V., Commereuc, S., Čermák, R.; Rheology upon photodegradation of polybutene-1. 42éme Colloque du Groupe Francais de Rhélogie, Clermont-Ferrand, France, 2007

11. Beníček, L., Verney, V., Commereuc, S., Obadal, M., Čermák, R.; Rheology evolution of polybutene-1 upon photodegradation; European Polymer Congress 2007, Portorož, Slovenia, 2007

12. Beníček, L., Obadal, M., Čermák, R.; Non-isothermal high-pressure crystallization of β-polypropylenes; European Polymer Congress 2007, Portorož, Slovenia, 2007

13. Beníček, L., Obadal, M.; High-pressure crystallization of polypropylenes; 10th International Research/Expert Conference „Trends in The Development of Machinery and Associated Technology“ TMT 2006, Lloret de Mar, Spain, 2006

14. Polášková, M., Čermák, R., Beníček, L., Obadal, M.; Polymer blends with microfibrillar-phase morphology; 10th International Research/Expert Conference „Trends in The Development of Machinery and Associated Technology“ TMT 2006, Lloret de Mar, Spain, 2006

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15. Obadal, M., Čermák, R., Beníček, L., Hnidáková, D.; Structure and properties of cast films produced from β-nucleated polypropylenes; 10th International Research/Expert Conference „Trends in The Development of Machinery and Associated Technology“ TMT 2006, Lloret de Mar, Spain, 2006

16. Čermák, R., Obadal, M., Habrová, V., Beníček, L.; Alpha and beta polypropylenes: The effects of processing on final properties; ANTEC 2006 - Society of Plastic Engineers, Charlotte NC, USA, 2006

17. Obadal, M., Čermák, R., Beníček, L.,Stoklasa, K.; Tailoring of multicomponent crystalline systems in isotactic polypropylene; 23rd discussion conference – Current and future trends in polymeric materials, 2005 Prague meetings on macromolecules, Prague, CZ, 2005

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CURRICULUM VITAE Personal data:

Name: Lubomír Beníček

Date of Birth: 1981, July 16

Present Address: Sadová 942, 687 51 Nivnice, Czech Republic

E-mail: Benicek@ft.utb.cz

Education:

2006 - 2009 Joint Ph.D. degree programme: - Laboratoire de Photochimie Moleculaire et macromoleculaire, Blaise Pascal Université in Clermont-Ferrand, France - Department of Polymer Engineering, Faculty of technology, Tomas Bata University in Zlín

2005 – 2006 Department of Polymer Engineering, Faculty of technology, Tomas Bata University in Zlín Ph.D. degree programme: Chemistry and Technology of Materials – Technology of Macromolecular Substances

1999 – 2005 Faculty of technology, Tomas Bata University in Zlín Master degree programme: Chemistry and Technology of Materials - Rubber and Plastic Technology Master Thesis: High-pressure crystallization of polypropylene (in English) appreciated by Dean Award for the best Master Thesis.

1995 – 1999 Technical College in Uherský Brod

Employment:

From 2009 Assistant - Department of Polymer Engineering, Faculty of technology, Tomas Bata University in Zlín

8 – 9 / 2005 Aliachem a.s., Fatra Napajedla – consultant to injection molding processing and technology

1999 - 2004 KASTEK komponenty s.r.o., – production of shoes components, technical and automotive parts

Others:

2 – 10 / 2006 Member of student chamber at Academic Senate of the Faculty of Technology; Student representative at Dean’s Council of the Faculty of technology

From 2005 Society of Plastics Engineers – member

RÉSUMÉ

Cette thèse porte sur l’inter relation entre la superstructure moléculaire, les propriétés et la dégradabilité du poly butène isotactique. L’évolution la plus importante est une transformation de phase particulière de la forme cristalline II à la forme cristalline I après cristallisation du polymère à partir de l’état fondu. Cette transformation a été caractérisée par calorimétrie, par diffraction des rayons X aux grands angles et par l’analyse des propriétés mécaniques en fonction des conditions de thermique de recuit. L’effet de la mise en forme sur la superstructure résultante a été mis en évidence sur des échantillons obtenus par injection ou par extrusion en utilisation les propriétés mécaniques en traction. Du fait de la très grande sensibilité du PB-1 vis-à-vis de la photodégradation des effets significatifs sur les propriétés thermiques et sur la transformation de phase ont été observés. Une caractérisation multi-échelle en photodégradation a été réalisée à partir d’analyses calorimétriques, d’analyse en spectroscopie infrarouge, de la diffraction des rayons X et des propriétés mécaniques en traction et ceci tant en photovieillissement en conditions accélérées qu’en conditions d’exposition naturelle.

ABSTRACT

This dissertation work is focused on the interrelations between supermolecular structure, properties and degradability of isotactic poly(1-butene) (PB-1). The most important evolution in PB-1 is a peculiar phase transformation from form II to form I after crystallization from melt. This phase transformation is characterized using calorimetry, wide-angle X-ray scattering and mechanical testing where changes in dependence on annealing conditions in time are observed. The effect of processing on resulting supermolecular structure is proven on injection-molded and extruded specimens using tensile testing. Because of high sensitivity of PB-1 to photodegradation, significant effects on thermal properties and the phase transformation is observed. Multiscale characterization is performed using calorimetric measurements, infra-red spectroscopy, rheology, X-ray scattering, microscopic techniques and tensile testing to describe consequences of photodegradation. As well, the effect of various degradation processes – natural weathering and accelerated UV irradiation lead to different results affecting resulting macroscopic properties.

ABSTRAKT

Tato dizertační práce je zaměřena na studium vzájemných vztahů mezi nadmolekulární strukturou, vlastnostmi a degradabilitou izotaktického poly(1-butenu) (PB-1). Nejvýznamnější vývoj v PB-1 je jeho zvláštní fázová transformace z formy II do formy I, jenž nastává po krystalizaci z taveniny. Tato fázová transformace je charakterizována pomocí kalorimetrického měření, široko-úhlé rentgenografie a mechanickými zkouškami, kde byly pozorovány změny v závislosti na podmínkách stárnutí. Vliv zpracovatelské technologie – vstřikování a vytlačování - na výslednou nadmolekulární strukturu byl prokázán na základě mechanických zkoušek. Jelikož je PB-1 velmi náchylný k fotodegradaci, významné změny byly zaznamenány v termálních vlastnostech a fázové transformaci. Široce zaměřená charakterizace důsledků fotodegradace byla provedena pomocí kalorimetrických měření, infračervené spektroskopie, reologie, rentgenografie, mikroskopických technik a tahových zkoušek. Rovněž byl posouzen vliv rozdílných degradačních procesů – venkovní stárnutí a zrychlené UV ozařování – které vedly k odlišným výsledkům ovlivňující výsledné makroskopické vlastnosti.